Hard magnetic alloy, hard magnetic alloy compact, and method for producing the same

ABSTRACT

A hard magnetic alloy in accordance with the present invention is composed of at least element T selected from the group consisting of Fe, Co and Ni, at least one rare earth element R, and boron (B). The hard magnetic alloy has an absolute value of the temperature coefficient of magnetization of 0.15%/° C. or less and a coercive force of 1 kOe, when being used in a shape causing a permeance factor of 2 or more. A hard magnetic alloy compact in accordance with the present invention has a texture, in which at least a part or all of the texture comprises an amorphous phase or fine crystalline phase having an average crystal grain size of 100 nm or less, is subjected to crystallization or grain growth under stress, such that a mixed phase composed of a soft magnetic or semi-hard magnetic phase and a hard magnetic phase is formed in the texture, and anisotropy is imparted to the crystal axis of the hard magnetic phase.

BACKGROUND OF THE INVENTION

1. Field of the Invention

The present invention relates to a hard magnetic alloy having excellentmagnetic characteristics and temperature-dependent properties and usedin sensors such as magnetic encoders and potentiometers, motors,actuators, and speakers. The present invention relates to a hardmagnetic alloy compact and a method for producing the alloy and thecompact.

2. Description of the Related Art

Nd-Fe-B magnets and Sm-Co magnets are generally known as magneticmaterials which show superior characteristics to that of ferrite magnetsand Alnico (Al-Ni-Co-Fe) magnets. Novel alloy magnets having furtherimproved characteristics and particularly Sm-Fe-N magnets have also beenintensively studied. These magnets, however, must contain at least 10atomic % of Nd or at least 8 atomic % of Sm. Use of large quantities ofexpensive rare earth elements inevitably increases the production costs.Since the magnetic characteristics of Nd-Fe-B magnets are largelydependent on temperature, they cannot be used as sensors. On the otherhand, Sm-Co magnets have not been used in practice in spite of theirsmaller thermal coefficients of magnetization, because they are moreexpensive than the Nd-Fe-B magnets.

Ferrite magnets and Alnico magnets are inexpensive compared to the rareearth magnets; however, the ferrite magnets have larger thermalcoefficients of magnetization and thus cannot be used as sensors,whereas the Alnico magnets have extremely low coercive forces.

The above-mentioned hard magnetic alloys have been produced by sprayingmolten alloys onto rotating drums to form thin ribbons by quenching thealloys or by spraying molten alloys into cooling gas to form alloypowders by quenching alloy droplets. The thin ribbons and alloy powdersmust therefore be formed into given shapes before being used for motors,actuators, and speakers.

Typical methods for molding magnetic powder include compaction moldingand injection compacting of a mixture of the magnetic powder and arubber or plastic binder. The resulting magnet is referred to as a “bondmagnet”. Since the versatility of possible form features of bond magnetsis high, they have been widely used in electronic parts. The binder inbond magnets, however, causes inferior magnetic characteristics becauseof decreased remanent magnetization and low mechanical strength of thebond magnet.

Accordingly, the advent of inexpensive magnetic materials having hardmagnetic characteristics superior to those of ferrite magnets andexcellent temperature-dependent properties has been eagerly awaited.

The present inventors have studied inexpensive hard magnetic materialshaving excellent magnetic characteristics and temperature-dependentproperties, and have discovered from various experimental results thatthe thermal coefficient of magnetization is related to the permeancefactor p.

Also, the present inventors have directed their attention to the heatingrate during annealing of a quenched alloy essentially consisting of anamorphous phase and discovered that hard magnetic characteristics arerelated to the nano-crystalline structure (particularly, crystal grainsize of the bcc(body centered cubic)-Fe phase) in a fine crystallinephase which is precipitated by the annealing.

SUMMARY OF THE INVENTION

It is an object of the present invention to provide a hard magneticalloy which is capable of low cost production and has excellent hardmagnetic characteristics and excellent temperature-dependent properties.

It is another object of the present invention to provide a hard magneticalloy compact having high mechanical strength and excellent magneticcharacteristics.

It is a further object of the present invention to provide a method forproducing the hard magnetic alloy or hard magnetic alloy compact.

A first aspect of the present invention is a hard magnetic alloycomprising at least one element T selected from the group consisting ofFe, Co and Ni, at least one rare earth element R, and B, the hardmagnetic alloy containing at least 10 percent by volume of a softmagnetic or semi-hard magnetic phase having a coercive force of 1 kOe orless and at least 10 percent by volume of a hard magnetic phase having acoercive force of 1 kOe or more, the absolute value of the thermalcoefficient of magnetization being 0.15%/° C. or less when the hardmagnetic alloy is used in a shape causing a permeance factor of 2 ormore.

Preferably, the hard magnetic alloy may primarily contain a finecrystalline phase having an average crystal grain size of 100 nm orless.

Preferably, the absolute value of the thermal coefficient ofmagnetization may be 0.1%/° C. or less when the hard magnetic alloy isused in a shape causing a permeance factor of 10 or more.

Preferably, the ratio Ir/Is of the remanent magnetization Ir to thesaturation magnetization Is may be 0.6 or more.

Preferably, the hard magnetic alloy may have the following formula:

T_(x)M_(y)R_(z)B_(w)

wherein T represents at least one element selected from the groupconsisting of Fe, Co and Ni, M represents at least one element selectedfrom the group consisting of Zr, Nb, Ta and Hf, R represents at leastone rare earth element, and the suffixes x, y, z and w by atomic percentsatisfy 50≦x, 0≦y≦15, 3≦z≦20, and 2≦w≦20, respectively. Preferably, thesuffixes x, y, z and w by atomic percent may satisfy 80≦x≦92, 1≦y≦5,3≦z≦10, and 3≦w≦7, respectively.

Preferably, the hard magnetic alloy may have the following formula:

T_(x)M_(y)R_(z)B_(w)Si_(u)

wherein T represents at least one element selected from the groupconsisting of Fe, Co and Ni, M represents at least one element selectedfrom the group consisting of Zr, Nb, Ta and Hf, R represents at leastone rare earth element, and the suffixes x, y, z, w, and u by atomicpercent satisfy 50≦x, 0≦y≦15, 3≦z≦20, 2≦w≦20, and 0≦u≦5, respectively.Preferably, the suffixes x, y, z, w, and u by atomic percent may satisfy80≦x≦92, 1≦y≦5, 3≦z≦10, 3≦w≦7, and 0.5≦u≦5, respectively.

Preferably, the hard magnetic alloy may have the following formula:

T_(x)M_(y)R_(z)B_(w)E_(v)

wherein T represents at least one element selected from the groupconsisting of Fe, Co and Ni, M represents at least one element selectedfrom the group consisting of Zr, Nb, Ta and Hf, R represents at leastone rare earth element, E represents at least one element selected fromthe group consisting of Cr, Al, Pt and platinum elements, and thesuffixes x, y, z, w, and v by atomic percent satisfy 50≦x, 0≦y≦15,3≦z≦20, 2≦w≦20, and 0≦v≦10, respectively. Preferably, the suffixes x, y,z, w, and v by atomic percent may satisfy 80≦x≦92, 1≦y≦5, 3≦z≦10, 3≦w≦7,and 0≦v≦5, respectively.

Preferably, the hard magnetic alloy may have the following formula:

T_(x)M_(y)R_(z)B_(w)E_(v)Si_(u)

wherein T represents at least one element selected from the groupconsisting of Fe, Co and Ni, M represents at least one element selectedfrom the group consisting of Zr, Nb, Ta and Hf, R represents at leastone rare earth element, E represents at least one element selected fromthe group consisting of Cr, Al, Pt and platinum elements, and thesuffixes x, y, z, w, v, and u by atomic percent satisfy 50≦x, 0≦y≦15,3≦z≦20, 2≦w≦20, 0≦v≦10, and 0≦u≦5, respectively. Preferably, thesuffixes x, y, z, w, v, and u by atomic percent satisfy 80≦x≦92, 1≦y≦5,3≦z≦10, 3≦w≦7, 0≦v≦5, and 0.5≦u≦5, respectively.

A second aspect of the present invention is a method for producing ahard magnetic alloy comprising the steps of: preparing an alloycontaining at least one element T selected from the group consisting ofFe, Co and Ni, at least one rare earth element R, and B, and essentiallyconsisting of an amorphous phase by a liquid quenching process, andannealing the alloy at a heating rate of 10° C./min. or more.

Preferably, a fine crystalline phase having an average crystal grainsize of 100 nm or less may be precipitated as a main phase by theannealing.

Preferably, the hard magnetic alloy in this method may have one of theabove-mentioned composition.

A third aspect of the present invention is a hard magnetic alloy compactcomprising an Fe-based or FeCo-based alloy containing 3 to 20 atomicpercent of at least one rare earth element R, and 2 to 20 atomic percentof B, wherein the alloy having a texture, in which at least a part orall of the texture comprises an amorphous phase or fine crystallinephase having an average crystal grain size of 100 nm or less, issubjected to crystallization or grain growth under stress, such that amixed phase composed of a soft magnetic or semi-hard magnetic phase anda hard magnetic phase is formed in the texture, anisotropy is impartedto the crystal axis of the hard magnetic phase, and the hard magneticalloy compact has a coercive force of 1 kOe or more.

Preferably, the hard magnetic alloy compact may comprise at least 10percent by volume of a soft magnetic or semi-hard magnetic phase havinga coercive force of 1 kOe or less which comprises a body centered cubic(bcc) Fe phase or bcc-FeCo phase, an Fe-B compound phase, and anamorphous phase as precipitates, and at least 10 percent by volume of ahard magnetic phase having a coercive force of 1 kOe or more whichcomprises an R₂Fe₁₄B phase, wherein R represents at least one rare earthelement.

Preferably, the annealed alloy may be crystallized or may be subjectedto crystal growth under stress and may be simultaneously compacted.

Preferably, the hard magnetic alloy contains an amorphous phase, and maybe formed by solidifying an alloy having hard magnetic characteristicswhen being crystallized, by means of a softening phenomenon of the alloyduring the crystallization reaction.

Preferably, the alloy may be heated under stress.

Preferably, the relative density of the compact obtained by compactingthe alloy is 90% or more.

Preferably, the hard magnetic alloy compact may have a remanentmagnetization of 100 emu/g or more.

Preferably, the ratio of the remanent magnetization Ir to the saturationmagnetization Is may be 0.6 or more.

Preferably, the hard magnetic alloy compact has one of the compositiondescribed in the first aspect.

A fourth aspect of the present invention is a method for producing ahard magnetic alloy compact comprising the following steps of: quenchingan Fe- or FeCo-based alloy containing 3 to 20 atomic percent of at leastone rare earth element R and 2 to 20 atomic percent of B so as to form atexture, in which at least a part or all of the texture comprises anamorphous phase or fine crystalline phase having an average crystalgrain size of 100 nm or less; performing crystallization or grain growthof the alloy under stress, such that a mixed phase composed of a softmagnetic or semi-hard magnetic phase and a hard magnetic phase is formedin the texture; imparting anisotropy to the crystal axis of the hardmagnetic phase.

Preferably, after performing crystallization or grain growth of thealloy under stress, the alloy may be annealed at 400° C. to 1,000° C. soas to precipitate a fine crystalline phase having an average crystalgrain size of 100 nm or less as a main phase in the texture.

Preferably, the alloy after quenching may be compacted while performingcrystallization or grain growth of the alloy under stress.

Preferably, the hard magnetic alloy contains an amorphous phase, and maybe formed by solidifying an alloy having hard magnetic characteristicswhen being crystallized, by means of a softening phenomenon of the alloyduring the crystallization reaction.

Preferably the alloy may be heated under stress.

BRIEF DESCRIPTION OF THE DRAWINGS

FIG. 1 is a cross-sectional view of the main section of a spark plasmasintering apparatus used in a method for producing a hard magnetic alloycompact in accordance with the present invention;

FIG. 2 is a graph of a waveform of a pulse sequence which is applied toan alloy powder in the spark plasma sintering apparatus shown in FIG. 1;

FIG. 3 is a front view of an entire spark plasma sintering apparatusused in a method for producing a hard magnetic alloy compact inaccordance with the present invention;

FIG. 4 is an isometric view of an embodiment using a hard magnetic alloyin accordance with the present invention as a magnet for a hallpotentiometer;

FIG. 5 is an isometric view of an embodiment using a hard magnetic alloyin accordance with the present invention as a rotary encoder magnet;

FIG. 6 is a cross-sectional view of a first embodiment using a hardmagnetic alloy in accordance with the present invention as a speakermagnet;

FIG. 7 is a cross-sectional view of a second embodiment using a hardmagnetic alloy in accordance with the present invention as a speakermagnet;

FIG. 8 is a graph of demagnetizing curves in the second quadrant of analloy ribbon having a composition of Fe₇₆Co₁₀Nb₂Pr₇B₅ at 29.5° C. to216° C.;

FIG. 9 is a graph of demagnetizing curves in the second quadrant of analloy ribbon having a composition of Fe₆₆Co₂₀Nb₂Pr₇B₅ at 35° C. to 198°C.;

FIG. 10 is a graph of demagnetizing curves in the second quadrant of analloy ribbon having a composition of Fe₈₄Nb₂Pr₇B₅Si₂ at 28.5° C. to 204°C.;

FIG. 11 is a graph illustrating the relationship between the magneticcharacteristics and the temperature of alloy ribbons in accordance withthe present invention and of a magnet of a comparative example;

FIG. 12 is a graph illustrating the dependence of Ir on temperature whenan alloy ribbon in accordance with the present invention and a magnet ofa comparative example are used at a shape causing p=1.5 or P=10;

FIG. 13 is a graph of the relationship between the permeance factor andthe temperature of a sintered alloy compact having a composition inaccordance with the present invention, of a ribbon alloy having acomposition in accordance with the present invention, and of aconventional Nd-Fe-B-based magnet;

FIG. 14 is a graph illustrating the dependence of magnetization ontemperature of two hard magnetic alloys having compositions ofFe₈₈Nb₂Pr₅B₅ and Fe₈₈Pr₇B₅, respectively;

FIG. 15 is a graph of demagnetizing curves in the second quadrant ofhard magnetic alloys having compositions of Fe₈₈Nb₂Pr₅B₅ and Fe₈₈Pr₇B₅;

FIG. 16 is a graph illustrating the dependence of the lattice constantand average crystal grain size on the heating rate of the bcc-Fe phasein an Fe₈₈Nb₂Pr₅B₅ alloy annealed at 700° C. to 750° C.;

FIG. 17 is a graph illustrating the dependence of the lattice constantand average crystal grain size on the heating rate of the bcc-Fe phasein an Fe₈₈Nb₂Nd₅B₅ alloy annealed at 700° C. to 750° C.;

FIG. 18 is a graph illustrating the dependence of the magneticcharacteristics on the heating rate in an Fe₈₈Nb₂Pr₅B₅ alloy which isannealed at 700° C. to 750° C. for 5 minutes and cooled at a coolingrate which is the same as the heating rate;

FIG. 19 is a graph illustrating dependence of the magneticcharacteristics on the heating rate in an Fe₈₈Nb₂Nd₅B₅ alloy which isannealed at 700° C. to 750° C. for 5 minutes and cooled at a coolingrate which is the same as the heating rate;

FIG. 20 is a graph illustrating the dependence of the magneticcharacteristics on the average crystal grain size of the bcc-Fe phase inFe₈₈Nb₂(Pr,Nd)₅B₅ alloys;

FIG. 21 is a graph illustrating the dependence of the magneticcharacteristics on the annealing time (holding time) in anFe₈₄Nb₂Nd₇B₅Si₂ alloy;

FIG. 22 is a graph of X-ray diffraction patterns at various annealingtimes (holding times) of an Fe₈₄Nb₂Nd₇B₅Si₂ alloy which is heated at aheating rate of 18° C./min. and annealed at 700° C.;

FIG. 23 is a graph of X-ray diffraction patterns at various annealingtimes (holding times) of an Fe₈₄Nb₂Nd₇B₅Si₂ alloy which is heated at aheating rate of 18° C./min. and annealed at 750° C.;

FIG. 24 is a graph of X-ray diffraction patterns at various annealingtimes (holding times) of an Fe₈₄Nb₂Nd₇B₅Si₂ alloy which is heated at aheating rate of 18° C./min. and annealed at 800° C.;

FIG. 25 is a graph illustrating the dependence of demagnetization curvesin the second quadrant on the heating rate in an Fe₈₈Nb₂Nd₅B₅ alloy,which is annealed at 750° C. for 180 seconds and then quenched;

FIG. 26 is a graph illustrating the dependence of magneticcharacteristics on the heating rate in an Fe₈₈Nb₂Nd₅B₅ alloy, which isannealed at 750° C. for 180 seconds and then quenched;

FIG. 27 is a graph illustrating the dependence of magneticcharacteristics on the annealing temperature (holding temperature) in anFe₈₈Nb₂Nd₅B₅ alloy, which is heated at a heating rate of 3 to 180°C./min., and annealed for 180 seconds and then quenched;

FIG. 28 is a graph illustrating the dependence of magnetization on thetemperature in an Fe₈₈Nb₂Nd₅B₅ alloy, which is heated at a heating rateof 3 to 180° C., and annealed at 750° C. for 180 seconds and thenquenched;

FIG. 29 is a graph illustrating the dependence of the crystal grain sizeon the heating rate in the bcc-Fe phase and the Nd₂Fe₁₄B phase in anFe₈₈Nb₂Nd₅B₅ alloy, which is annealed at 750° C. for 180 seconds andthen quenched;

FIG. 30 is differential scanning calorimetric (DSC) thermograms of anamorphous alloy ribbon having a composition of Fe₈₈Nb₂Nd₅B₅, which isobtained by a quenching process;

FIG. 31A is a graph illustrating crystallization temperatures T_(x1) andT_(x2), respectively, of the bcc-Fe phase and Nd₂Fe₁₄B phase at variousheating rates, which are determined by the DSC thermograms shown in FIG.30; and

FIG. 31B is a graph illustrating the difference (D_(tx)=T_(x2)−T_(x1))between the crystallization temperatures T_(x1) and T_(x2),respectively, of the bcc-Fe phase and Nd₂Fe₁₄B phase at various heatingrates, which are determined by the DSC thermograms shown in FIG. 30;

FIG. 32 is an isometric view illustrating the direction of the sinteringpressure applied during production of a compact;

FIG. 33A includes TMA curves (A) and FIG. 33B a DSC thermogram (B) of anamorphous alloy in accordance with the present invention and of acrystalline alloy not containing an amorphous component;

FIGS. 34A to 34C are photographs by microscopy of a compact obtained byvarying the sintering temperature of an amorphous alloy in accordancewith the present invention;

FIG. 35 is a graph of X-ray diffraction patterns of a compact obtainedby various sintering temperatures of an amorphous alloy in accordancewith the present invention;

FIG. 36 is a graph illustrating the relationship between the density andthe sintering temperature in a compact in accordance with the presentinvention and a compact in a comparative example;

FIG. 37 is a graph illustrating the relationship between the density andthe sintering pressure in a compact in accordance with the presentinvention and a compact in a comparative example;

FIG. 38 is a graph of X-ray diffraction patterns of an amorphous alloypowder in accordance with the present invention;

FIG. 39 is a graph of X-ray diffraction patterns of an amorphous alloypowder in accordance with the present invention;

FIG. 40 is a graph illustrating the relationship between the sinteringtime and the temperature and between the sintering time and theexpansion when sintering amorphous Fe₈₈Nb₂Nd₅B₅ and Fe₈₆Nb₂Nd₇B₅powders, and a nano-crystalline Fe₈₈Nb₂Nd₅B₅ powder;

FIG. 41 is a graph illustrating the relationship between the compacteddensity and the sintering temperature of compacts obtained by molding anamorphous Fe₈₈Nb₂Nd₅B₅ powder, an amorphous Fe₈₆Nb₂Nd₇B₅ powder, and acrystalline Fe₈₈Nb₂Nd₅B₅ powder;

FIG. 42 is a graph of X-ray diffraction patterns of an Fe₈₆Nb₂Nd₇B₅compact and an Fe₈₈Nb₂Nd₅B₅ compact immediately after sintering;

FIG. 43 is a graph of the relationship between the magneticcharacteristics and the sintering temperature when sintering anamorphous Fe₈₈Nb₂Nd₅B₅ powder and an amorphous Fe₈₆Nb₂Nd₇B₅ powder undera pressure of 636 MPa;

FIG. 44 is a graph of the relationship between the magneticcharacteristics and the sintering temperature when sintering anamorphous Fe₈₈Nb₂Nd₅B₅ powder under a pressure of 636 MPa;

FIG. 45 is a graph of the relationship between the magneticcharacteristics and the sintering pressure when sintering an amorphousFe₈₈Nb₂Nd₅B₅ powder at a temperature of 600° C.;

FIG. 46 is a graph of the relationship between the magneticcharacteristics and the density of a compact when molding an amorphousFe₈₈Nb₂Nd₅B₅ powder;

FIG. 47 is a graph of X-ray diffraction patterns after annealing at 750°C. of an Fe₈₈Nb₂Nd₅B₅ compact and an Fe₈₆Nb₂Nd₇B₅ compact which aresintered under a pressure of 636 MPa at a sintering temperature of 600°C.;

FIG. 48 is a graph of magnetization curves after annealing at 750° C. ofan Fe₈₈Nb₂Nd₅B₅ compact and an Fe₈₆Nb₂Nd₇B₅ compact which are moldedunder a pressure of 636 MPa at a sintering temperature of 600° C.;

FIG. 49 is a graph of magnetization curves after annealing at 750° C. ofan amorphous Fe₈₈Nb₂Nd₅B₅ powder as an example and a crystallineFe₈₈Nb₂Nd₅B₅ powder as a comparative example which are molded under apressure of 636 MPa at a heating rate of 1.8° C./sec., a sinteringtemperature of 600° C., and a holding time of 480 sec. (8 min.);

FIG. 50 is a graph illustrating the magnetic characteristics afterannealing at 627° C. to 827° C. of a compact which is obtained bymolding an amorphous Fe₈₈Nb₂Nd₅B₅ powder under a pressure of 636 MPa ata sintering temperature of 500° C. to 600° C.;

FIG. 51 is a graph illustrating the magnetic characteristics afterannealing at 627° C. to 827° C. of a compact which is obtained bymolding an amorphous Fe₈₆Nb₂Nd₇B₅ powder under a pressure of 636 MPa ata sintering temperature of 500° C. to 600° C.;

FIG. 52 is a graph illustrating the demagnetization curves of a compactwhich is obtained by molding an amorphous Fe₉₀Nb₂Nd₅B₃ powder under apressure of 636 MPa at a sintering temperature of 600° C. for 8 min.;

FIG. 53 is a graph illustrating the demagnetization curves of a compactwhich is obtained by molding an amorphous Fe₈₉Nb₂Nd₄B₅ powder under apressure of 636 MPa at a sintering temperature of 600° C. for 8 min.;

FIG. 54 is a graph illustrating the demagnetization curves of a compactwhich is obtained by molding an amorphous Fe₇₆Co₁₀Nb₂Nd₇B₅ powder undera pressure of 636 MPa at a sintering temperature of 600° C. for 8 min.;

FIG. 55 is a graph illustrating the demagnetization curves of a compactwhich is obtained by molding an amorphous Fe₈₄Nb₂Nd₇B₅Si₂ powder under apressure of 636 MPa at a sintering temperature of 600° C. for 8 min.;

FIG. 56 is a graph illustrating a heating pattern;

FIG. 57 is a graph illustrating a heating pattern;

FIG. 58 is a graph of DSC thermograms at various heating rates of anamorphous alloy ribbon having a composition of Fe₈₈Nb₂Nd₅B₅ which isobtained by a quenching process;

FIG. 59 is a graph illustrating the dependence of magneticcharacteristics on T₁ in an Fe₈₈Nb₂Nd₅B₅ composition; and

FIG. 60 is a graph illustrating the dependence of magneticcharacteristics on T₂ in an Fe₈₈Nb₂Nd₅B₅ composition.

DESCRIPTION OF THE PREFERRED EMBODIMENTS

The preferred embodiments of the present invention will now beillustrated with reference to the drawings.

The hard magnetic alloy in accordance with the present inventioncomprises at least one element selected from the group consisting of Fe,Co and Ni, at least one rare earth element R, and B (boron), theabsolute value of the temperature coefficient of magnetization of thealloy is 0.15%/° C. or less when it is used in a shape causing apermeance factor of 2 or more, and its coercive force is 1 kOe or more.

Magnetic characteristics of magnetic materials are generally expressedby the hysteresis curves in the second quadrant, that is,demagnetization curves. Since magnetic materials after magnetization lieunder a demagnetizing field, which is a magnetic field in the reversedirection caused by remanent magnetization of itself, the operatingpoint (the magnetic flux density B of the material and the demagnetizingfield H) is given at a point p on the demagnetizing curve, wherein theB/μ₀H value, a nondimensional parameter, is referred to as a permeancefactor p, and a line OP between p and the origin O is referred to as apermeance line. The permeance factor p or permeance line depends on theshape of the magnet, decreases as the length of the magnetizationdirection decreases, and increases as the length increases, for example,p=1.5 for a disk magnet or p=10 for a prismatic magnet.

The following equation (I) stands between the permeance factor p and thedemagnetizing factor N:

p=(1−N)/N  (I)

The operating point (B, H) is therefore determined by thedemagnetization curve and the shape of the magnetic material. Themagnetostatic energy U generated in an outer field by the magneticmaterial is represented by the following equation (II):

U=BHV/2  (II)

wherein V is the volume of the magnetic material. The demagnetizationfield, that is, the permeance line changes with the shape of themagnetic material, hence the operating point p and thus themagnetostatic energy U change. The magnetostatic energy U has a maximumvalue at a given operating point p_(m), and the corresponding product ofthe magnetic flux density B and the demagnetizing field H is referred toas the maximum energy product (BH)_(max).

When using the hard magnetic material in accordance with the presentinvention in sensors, it is preferable that the hard magnetic materialhas an excellent temperature-dependent property, that is, a smallabsolute value of the temperature coefficient of magnetization, whichprevents the drift of the output caused by a change in temperature.Since the hard magnetic material in accordance with the presentinvention has a small absolute value of the temperature coefficient ofmagnetization of 0.15%/° C. or less when being used in a shape causing apermeance factor of 2 or more as described above, it can be preferablyused for sensors. It is more preferable that the hard magnetic alloy inaccordance with the present invention be used in a shape causing apermeance factor of 10 or more, in order to achieve a further improvedtemperature-dependent property, that is, a smaller absolute value of thetemperature coefficient of magnetization of 0.1%/° C. or less. Such alow absolute value of the temperature coefficient of magnetization inthe hard magnetic alloy in accordance with the present invention iscomparable to or superior to any conventional Nd-Fe-B-based magnetshaving absolute values of temperature coefficients of magnetization in arange from 0.11 to 0.15%/° C. when being used in a shape causing apermeance factor of 10 or more.

Further, the hard magnetic alloy in accordance with the presentinvention has a larger coercive force compared to Alnico magnets, and ismore inexpensive than conventional Sm-Co-based magnets having excellenttemperature-dependent properties.

In the hard magnetic alloy in accordance with the present invention, Si(silicon) is preferably substituted for 0.5 to 5 atomic percent of theelement T, or 0.5 to 20 atomic percent of Co (cobalt) is preferablycontained in the element T, in order to further improvetemperature-dependent properties.

The hard magnetic alloy in accordance with the present invention maycontain at least 10 percent by volume of a soft magnetic or semi-hardmagnetic phase having a coercive force of 1 kOe or less and at least 10percent by volume of a hard magnetic phase having a coercive force of 1kOe or more, and may have a coercive force of 1 kOe or more. The hardmagnetic alloy with such a microstructure has an intermediatecharacteristic between the soft magnetic or semi-hard magnetic phase andthe hard magnetic phase. Since the soft magnetic phase generally hashigh magnetization, it improves the overall remanent magnetization ofthe hard magnetic alloy. Accordingly, the inclusion of the soft magneticphase is preferable for enhancing magnetic characteristics of the hardmagnetic alloy. When the content of the soft magnetic phase with acoercive force of 1 kOe or less is less than 10 percent by volume, alarge amount of Nd (neodymium) or the like, which is essential for thehard magnetic phase, must be added, and the remanent magnetizationdecreases. When the content of the hard magnetic phase with a coerciveforce of 1 kOe or more is less than 10 percent by volume, the coerciveforce of the hard magnetic alloy decreases. The preferred content of thesoft magnetic phase with a coercive force of 1 kOe or less is 20 to 60percent by volume, and the preferred content of the hard magnetic phasewith a coercive force of 1 kOe or more is 40 to 80 percent by volume.

It is preferable that the hard magnetic alloy in accordance with thepresent invention contains at least 10 percent by volume of a magneticphase with a Curie point of 600° C. or more and at least 10 percent byvolume of a magnetic phase with a Curie point of 600° C. or less, andhas a coercive force of 1 kOe or more, such that the hard magnetic alloyhas an intermediate characteristic between the soft magnetic phase andthe hard magnetic phase. The bcc Fe phase has a Curie point of near 770°C. and the R₂Fe₁₄B phase has a Curie point of near 315° C. The hardmagnetic alloy in accordance with the present invention must contain themagnetic phase with a Curie point of 600° C. or more and the magneticphase with a Curie point of 600° C. or less, in order that the hardmagnetic alloy contains the soft magnetic phase and the hard magneticphase which participates in magnetization.

If the content of the magnetic phase with a Curie point of 600° C. ormore is less than 10 percent by volume, the magnetization significantlychanges with temperature in use at a relatively high permeance. On theother hand, if the content of the magnetic phase with a Curie point of600° C. or less is less than 10 percent by volume, the coercive force ofthe hard magnetic alloy decreases because of a relatively small contentof hard magnetic phase. The preferable content of the magnetic phasewith a Curie point at 600° C. or more is 20 to 60 percent by volume, andthe preferable content of the magnetic phase with a Curie point of 600°C. or less is 40 to 80 percent by volume.

The hard magnetic alloy in accordance with the present inventionprimarily contains a fine crystalline phase with an average crystalgrain size of 100 nm or less as a precipitate. The fine crystallineprecipitate is composed of a bcc-Fe phase with an average size of 100 nmor less which primarily forms the soft magnetic phase and a R₂Fe₁₄Bphase with an average size of 100 nm or less which primarily forms thehard magnetic phase. Further, the hard magnetic alloy in accordance withthe present invention has a nano-composite texture including the finecrystalline phase composed of the bcc-Fe phase and R₂Fe₁₄B phase, and aresidual amorphous phase.

The hard magnetic alloy in accordance with the present invention may beproduced by quenching the alloy melt having the above-mentionedcomposition to prepare an alloy primarily containing an amorphous phase,that is, amorphous alloy, and by annealing the amorphous alloy.

It is preferable that the hard magnetic alloy in accordance with thepresent invention has a coercive force of 2 kOe or more.

It is preferable that the hard magnetic alloy in accordance with thepresent invention has a ratio Ir/Is of the remanent magnetization Ir tothe saturation magnetization Is of 0.6 or more.

The average crystal grain size of the crystalline phase in the hardmagnetic alloy and the concentration of each element in each phase canbe controlled by the annealing conditions of the amorphous alloy, suchas the heating rate, the annealing temperature, and the holding time.

The hard magnetic alloy in accordance with the present invention isproduced as follows.

The processes preparing the amorphous alloy include the formation of anamorphous alloy ribbon by spraying an alloy melt onto a drum to quenchthe alloy melt; the formation of an alloy powder by spraying droplets ofan alloy melt into a cooling gas to quench the alloy melt; and asputtering or CVD process.

The annealing of the amorphous alloy can be performed with any heatingmeans. For example, in a process for preparing a hard magnetic alloycompact, the amorphous alloy is pulverized, and the pulverized amorphousalloy is pressure-molded with a hot press, and is simultaneouslyannealed at an adequate heating rate and annealing temperature. Thepreferable heating rate depends on the composition of the amorphousalloy, and is generally 10° C./min. or more, and specifically 100°C./min. or more. A heating rate of less than 10° C./min. causescoarsening of crystal grains which are precipitated in the alloy duringthe heating process, hence exchange coupling characteristics between thebcc-Fe soft magnetic phase and the R₂Fe₁₄B hard magnetic phase isdeteriorate, that is, the hard magnetic characteristics deteriorate. Ifthe heating rate is 100° C./min. or more, a more homogeneous finetexture can be formed. The upper limit of the heating rate by thecurrent apparatuses is near 200° C./min.

The annealing temperature is preferably 600° C. to 900° C., and morepreferably 700° C. to 750° C. The holding time (annealing time) ispreferably 1 min. to 20 min., and more preferably 3 m. to 10 min. Thecombination of the preferable annealing temperature and annealing timedepends on the composition of the amorphous alloy. At an annealingtemperature of less than 600° C., a relatively small amount of R₂Fe₁₄Bphase having hard magnetic characteristics is precipitated, and thus thehard magnetic alloy does not have satisfactory hard magneticcharacteristics. On the other hand, at an annealing temperature above900° C., other precipitates will form or crystal grains will becoarsened, resulting in deterioration of hard magnetic characteristics.

The hard magnetic alloy compact in accordance with the present inventionis produced as follows.

The hard magnetic alloy compact in accordance with the present inventionis basically composed of an Fe-based or FeCo-based alloy which contains3 to 20 atomic percent of at least one rare earth element R, and 2 to 20atomic percent of B. The alloy has a texture, in which at least a partor all of the texture is composed of an amorphous phase or finecrystalline phase having an average crystal grain size of 100 nm orless, is subjected to crystallization or grain growth under stress, suchthat a mixed phase composed of a soft magnetic or semi-hard magneticphase and a hard magnetic phase is formed in the texture, and anisotropyis imparted to the crystal axis of the hard magnetic phase. The Fe- orFeCo-based alloy is an alloy containing an amorphous phase, that is,amorphous alloy, and the amorphous alloy may contain a small amount ofcrystalline phase and shows hard magnetic characteristics when beingcrystallized.

In the production of the hard magnetic alloy compact, an alloy powder isprepared. The process for preparing the alloy powder generally includesthe step for quenching an amorphous alloy melt to form an alloy ribbonor powder, and the step for pulverizing the alloy ribbon. The alloypowder has a particle size of approximately 50 μm to 150 μm. It ispreferable that the alloy powder is composed of only the amorphousphase.

The processes preparing the amorphous alloy, which may contain a smallamount of crystalline phase as described above, include the formation ofan amorphous alloy ribbon by spraying an alloy melt onto a drum toquench the alloy melt; the formation of an alloy powder by sprayingdroplets of an alloy melt into a cooling gas to quench the alloy melt;and a sputtering or CVD process. The resulting alloy ribbon or powdercontains a texture composed of a fine crystalline phase precipitate withan average crystal grain size of 100 nm or less or an amorphous phase.

Under stress, the amorphous phase in the alloy powder is crystallized orthe crystal grains in the fine crystalline phase are grown, and then thealloy powder is compacted. A mixed phase composed of a soft magnetic orsemi-hard magnetic phase and a hard magnetic phase is formed in thetexture, which is composed of the fine crystalline phase precipitatewith an average crystal grain size of 100 nm or less, or a finecrystalline phase having an average crystal grain size of 100 nm or lessis precipitated in the amorphous phase and the mixed phase issimultaneously formed. Further, anisotropy is imparted to the crystalaxis of the hard magnetic phase during the compaction process. Such ananisotropic crystal axis of the hard magnetic phase shows a higherremanent magnetization Ir than the isotropic crystal axis.

It is preferable that the amorphous alloy powder be heated when it issubjected to crystallization or grain growth under stress.

It is preferable that the amorphous alloy powder be compacted by meansof the softening of the amorphous alloy powder which accompanies thecrystallization. The amorphous alloy powder is noticeably softened inthe heating process for the crystallization, and the amorphous alloypowder particles are tightly bonded to each other under the pressure.The resulting hard magnetic compact therefore has a high density. It ispreferable that the amorphous alloy powder contains at least 50 percentby weight of amorphous phase in order to achieve tight bonding aftersolidification with heat and pressure and to obtain a permanent magnetwith strong hard magnetic characteristics.

Examples of the production of the compact from the alloy powder includecompaction of the amorphous alloy powder in a spark plasma sinteringsystem under pressure while applying pulse currents to the alloy powderto heat the amorphous alloy at near the crystallization temperature fora given time; or compaction of the amorphous alloy powder by the heatformed by the pulse currents applied to the alloy powder and by thepressure formed by pressing the alloy powder with two punches at nearthe crystallization temperature.

FIG. 1 is a cross-sectional view of the main section of a spark plasmasintering system which is preferably used for producing the hardmagnetic alloy compact in accordance with the present invention. Thespark plasma sintering system includes a tungsten carbide dice 1,tungsten carbide upper and lower punches 2 and 3 inserted in the dice 1,a tungsten carbide overcap dice 8 provided at the exterior of the dice1, a base table 4 which supports the lower punch 3 and functions as anelectrode for leading pulse currents, a base table 6 which pressesdownward the upper punch 2 and functions as another electrode forleading the pulse currents, a thermocouple 7 for measuring thetemperature of the alloy powder placed between the upper and lowerpunches 2 and 3.

FIG. 3 is a front view of the entire spark plasma sintering apparatus A,Model SPS-2050 made by Sumitomo Coal Mining Co., Ltd., in which thestructure of the main section is shown in FIG. 1.

The spark plasma sintering apparatus includes an upper plate 11 and alower plate 12, and a chamber 13 being in contact with the upper plate11. The main section shown in FIG. 1 is stored in the chamber 13. Thechamber 13 is connected to an evacuation system and an atmospheric gassupply system not shown in the drawing, so that the alloy powder 6placed between the upper and lower punches 2 and 3 is held in a givenatmosphere, such as an inert gas atmosphere. The upper and lower punches2 and 3 and base tables 4 and 5 are connected to an electrical powerunit not shown in FIGS. 1 and 3, which supplies a pulse sequence ofelectric current as shown in FIG. 2 to the alloy powder through theupper and lower punches 2 and 3 and base tables 4 and 5.

In the production of the compact using the spark plasma sintering systemshown in FIGS. 1 and 3, the alloy powder 6 is placed between the punches2 and 3, the chamber is evacuated, and the alloy powder 6 is molded bythe pressure added to the upper and lower punches 2 and 3, whileapplying the pulse currents to the alloy powder. The amorphous alloypowder is heated at or near the melting point for a given time understress, such that the amorphous alloy is crystallized or the crystalgrains grow. The hard magnetic alloy compact in accordance with thepresent invention is produced, for example, in such a manner.

The pressure applied during the spark plasma sintering process forcrystallization or grain growth ranges generally from 200 to 1,500 MPa,and preferably from 500 to 1,000 MPa. Satisfactory anisotropy in thehard magnetic phase is not achieved under a pressure of less than 200MPa. Further, the density of the resulting compact is low due to highporosity in the compact. On the other hand, a pressure higher than 1,500MPa causes damage of the tungsten carbide dice at the temperature.

The alloy powder 6 is heated at a heating rate of 10° C./min. or more,and preferably 20° C./min. or more. Crystal grains are coarsened at aheating rate of less than 10° C./min., and thus the hard magneticcharacteristics deteriorate because of decreased exchange coupling.

In the spark plasma sintering process, it is preferable that theamorphous alloy powder be sintered at a sintering temperature Tssatisfying the equation, Tx−200° C.≦Ts ° C.≦Tx+200° C., wherein Tx isthe starting temperature of the crystallization of the amorphous alloy.A satisfactory compact having a high density is not obtained at atemperature of less than Tx−200° C., whereas the hard magneticcharacteristics deteriorate because of the growth of the finecrystalline phase at a temperature of higher than Tx+200° C.

In the spark plasma sintering process using such a spark plasma system,the current flow enables rapid heating of the alloy powder at a givenrate and exact control of the temperature of the alloy powder comparedwith heating using a heater. The alloy powder can therefore be heatedaccording to a designed heating program.

In the above-mentioned process, the compact is formed by solidifying thealloy powder while or after the alloy powder is subjected tocrystallization or crystal grain growth under pressure by the sparkplasma sintering process. The alloy powder may be placed into a mold andmay be heated to or near the crystallization temperature of theamorphous alloy while being pressed by a hot press. The solidificationof the alloy powder, the crystallization and the grain growth for theproduction of the compact simultaneously proceed also in this process.

In the solidification process by softening of the alloy powder, thepressure, temperature and molding time are controlled such that theresulting compact has a relative density of 90% or more. Since thesintered compact has a significantly dense texture and high mechanicalstrength, it functions as a compact permanent magnet having strong hardmagnetic characteristics.

After the crystallization or grain growth under pressure, the alloypowder is annealed at 400° C. to 1,000° C. while or after the alloypowder is compacted. A fine crystalline phase with an average crystalgrain size of 100 nm or less is thereby precipitated as a main phase,resulting in the occurrence of hard magnetic characteristics. At anannealing temperature of less than 400° C., a relatively small amount ofR₂Fe₁₄B phase having hard magnetic characteristics is precipitated,hence the resulting compact does not have satisfactory hard magneticcharacteristics. At a temperature higher than 1,000° C., the crystalgrains in the fine crystalline phase grow, and thus the hard magneticcharacteristics deteriorate.

In particular, a compact having extremely high hard magneticcharacteristics is obtained under the following conditions: A finecrystalline phase having an average crystal grain size of 100 nm or lessoccupies 60 percent by volume of the compact and the balance is theamorphous phase; and a bcc-Fe phase or bcc-FeCo phase, and an R₂Fe₁₄Bphase, wherein R is at least one rare earth element, are formed in thefine crystalline phase.

In a preferred embodiment, a hard magnetic alloy compact havingintermediate characteristics between the soft magnetic phase and thehard magnetic phase can be provided under the following conditions: Thecompact contains at least 10 percent by volume of a soft or semi-hardmagnetic phase with a coercive force of 1 kOe or less and at least 10percent by volume of a hard magnetic phase with a coercive force of 1kOe or more; a bcc-Fe or bcc-FeCo phase, an Fe-B compound, and anamorphous phase are precipitated in the soft or semi-hard magnetic phasewith a coercive force of 1 kOe or less; and only an R₂Fe₁₄B phase,wherein R is at least one rare earth element, is precipitated in thehard magnetic phase with a coercive force of 1 kOe or more. Theformation of the soft magnetic phase having high magnetization in thecompact results in a significant increase in remanent magnetization ofthe compact. When the content of the soft magnetic phase with a coerciveforce of 1 kOe or less is less than 10 percent by volume, the content ofthe rare earth element, which is necessary for the hard magnetic phase,should be increased, although the coercive force of the compact isincreased. When the content of the hard magnetic phase with a coerciveforce of 1 kOe or more is less than 10 percent by volume, the content ofthe rare earth element, the coercive force of the compact is decreased.

The hard magnetic alloy compact prepared by the above-mentioned processhas exchange coupling characteristics between the fine soft magneticphase and the fine hard magnetic phase in the fine texture. Further, thehard magnetic alloy compact has a higher Fe content than that inconventional rare earth magnets, hence it can be used as a permanentmagnet having high magnetic characteristics, that is, a remanentmagnetization of 100 emu/g or more, and a remanence ratio Ir/Is of theremanent magnetization Ir to the saturation magnetization Is of 0.6 ormore. Since the hard magnetic alloy compact is subjected tocrystallization or crystal grain growth under stress, anisotropy isimparted to the crystal axis of the hard magnetic phase. The hardmagnetic alloy compact has high uniaxial anisotropy and thus a highremanent magnetization Ir. Further, the hard magnetic alloy compact isproduced by molding of the amorphous alloy powder under pressure, henceit has higher mechanical strength than conventional bond magnets.Accordingly, the hard magnetic alloy compact is used as a compactpermanent magnet having high hard magnetic characteristics. Further, ahard magnetic alloy compact having a given shape can be easily formedfrom the alloy powder.

Accordingly, the hard magnetic alloy compact in accordance with thepresent invention is useful as a permanent magnet which is used invarious apparatuses, for example, motors, actuators, speakers, and thematerial costs of these apparatuses can be reduced.

The preferred hard magnetic alloy in accordance with the presentinvention is represented by the following formula:

T_(x)M_(y)R_(z)B_(w)

wherein T is at least one element selected from the group consisting ofFe, Co and Ni. The content of the element T must be at least 50 atomicpercent as the main component for achieving hard magneticcharacteristics. The saturation magnetization Is increases as thecontent x of the element T increases. In order to achieve a highremanent magnetization Ir of 100 emu/g or more, a saturationmagnetization Is of at least 130 emu/g is required, and thus it ispreferable that the content of the element T be at least 80 atomicpercent. The hard magnetic alloy in accordance with the presentinvention must contain Fe among the elements T.

In the above-mentioned formula, M is at least one element selected fromthe group consisting of Zr, Nb, Ta, and Hf. The element M has a highability for forming an amorphous phase. The addition of the element M inthe hard magnetic alloy in accordance with the present invention enableseasy formation of the amorphous phase even when the content of the rareearth element R is low. As the content y of the element Y as asubstituent of the rare earth element R increases, the remanentmagnetization Ir increases whereas the coercive force iHc decreases,that is, the hard magnetic characteristics of the hard magnetic alloy ischanged to soft magnetic characteristics. As the content of theamorphous-forming element M as the substituent of the magnetic element Tincreases, the saturation magnetization Is and the remanentmagnetization Ir decrease. Accordingly, it is preferable for achievingsatisfactory magnetic characteristics that the concentration of theelement M be in a range from 0 atomic percent to 15 atomic percent, andmore preferably 1 atomic percent to 5 atomic percent.

In the above-mentioned formula, R is at least one rare earth element,such as Sc, Y, La, Ce, Pr, Nd, Pm, Sm, Eu, Gd, Tb, Dy, Ho, Er, Tm, Yband Lu. The intermetallic compound R₂Fe₁₄B, which is precipitated whenheating an amorphous alloy containing R, Fe and B to an appropriatetemperature in a range from 600° C. to 900° C., imparts satisfactoryhard magnetic characteristics to the hard magnetic alloy in accordancewith the present invention. As the content z of the element R increases,the saturation magnetization Ir decreases. In order to achieve a highremanent magnetization Ir of 100 emu/g or more, a saturationmagnetization Is of at least 130 emu/g is required, and thus it ispreferable that the content z of the element R be 20 atomic percent orless. When the content of the element R contributing to the formation ofthe amorphous phase is decreased, no satisfactory amorphous or finecrystalline phase is formed, hence the content of the element R ispreferably at least 2 atomic percent. When part or all of the element Ris composed of Nd and/or Pr, higher hard magnetic characteristics can beachieved.

In the above-mentioned formula, B (boron) is an element easily formingan amorphous phase. The compound R₂Fe₁₄B precipitated by annealing ofthe amorphous phase containing Fe and B at an appropriate temperature ina range from 600° C. to 900° C. imparts hard magnetic characteristics tothe hard magnetic alloy in accordance with the present invention. It ispreferable that the content of B be at least 2 atomic percent forobtaining a satisfactory amorphous or fine crystalline phase, however,the saturation magnetization Is, remanent magnetization Ir, and coerciveforce iHc decrease as the content w of B increases, hence the preferableupper limit of the B content is 20 atomic percent, and more preferably 7atomic percent.

The hard magnetic alloy in accordance with the present invention maycontain at least one element E selected from the group consisting of Cr,Al, Pt, Cu, Ag, Au, and platinum elements for fining the crystal textureand improving corrosion resistance of the alloy. In this case, the hardmagnetic alloy is represented by the following formula:

T_(x)M_(y)R_(z)B_(w)E_(v)

The content x of the magnetic element is preferably 50 atomic percent ormore, and more preferably 80 atomic percent or more and 92 atomicpercent or less in order to increase the saturation magnetization Is. Inparticular, it is preferable that the content x be 86 atomic percent ormore and 93 atomic percent or less for achieving a high remanentmagnetization Ir of 100 emu/g or more.

The content y of the element M is preferably 0 atomic percent or moreand 15 atomic percent or less, and more preferably 1 atomic percent ormore and 5 atomic percent or less in order to achieve excellent hardmagnetic characteristics. In particular, it is preferable that thecontent y be 0.5 atomic percent or more and 3 atomic percent or less forachieving a high remanent magnetization Ir of 100 emu/g or more.

The content z of the element R is preferably 3 atomic percent or moreand 20 atomic percent or less, and more preferably 3 atomic percent orless and 10 atomic percent or less in order to impart excellent hardmagnetic characteristics to the alloy and to form a satisfactoryamorphous or fine crystalline phase. In particular, it is preferablethat the content z be 3 atomic percent or more and 7 atomic percent orless for achieving a high remanent magnetization Ir of 100 emu/g ormore.

The content w of the element B is preferably 2 atomic percent or more inorder to form a satisfactory amorphous or fine crystalline phase.Further, it is preferable that the content w be 20 atomic percent orless, and more preferably 7 atomic percent or less for achievingexcellent hard magnetic characteristics. The addition of the element Ecauses the improved corrosion resistance of the hard magnetic alloy. Thehard magnetic characteristics, however, deteriorate if an excessiveamount of element E is present, hence the content w of the element E ispreferably 10 atomic percent or less, and more preferably 5 atomicpercent or less. When a high remanent magnetization Ir of more than 100emu/g is required, it is preferable that the element E be not added.

It is preferable that the hard magnetic alloy in accordance with thepresent invention contain Co as the element T other than Fe in order toreduce the absolute value of the temperature coefficient ofmagnetization when the alloy is used in a shape causing a permeancefactor of 2 or more, and particularly 10 or more. Since Co as theelement T in the alloy causes an increase in the Curie point, dependenceof the magnetization and coercive force on temperature decreases,dependence of the magnetic characteristics on temperature also decreasesbecause of an increased remanence ratio. Further, Co is contained in thebcc-Fe phase, hence dependence of the remanent magnetization ontemperature decreases. Since an excessive amount of Co in the alloycauses deterioration of magnetic characteristics, the content isdetermined in view of the alloy composition and annealing conditions.The Co content is generally 50 atomic percent or less, preferably 0.5atomic percent or more and 30 atomic percent or less, and morepreferably 0.5 atomic percent or more and 20 atomic percent or less.

When Si is added as a substituent of the element T in the hard magneticalloy in accordance with the present invention, magneticcharacteristics, such as coercive force iHc and maximum magnetic energyproduct (BH)_(max), can be further improved. Further, the addition of Sican decrease the absolute value of the temperature coefficient ofmagnetization when the alloy is used in a shape causing a permeancefactor of 2 or more, and particularly 10 or more. An excessive Sicontent causes an unintended decrease in hard magnetic characteristicsdue to a relatively low content of the element T. Accordingly, the Sicontent is preferably 0.5 atomic percent or more and 5 atomic percent orless, and more preferably 0.5 atomic percent or more and 3 atomicpercent or less, and the Si content is determined in view of the alloycomposition and annealing conditions.

The hard magnetic alloy having such improved coercive force andtemperature dependent properties can be suitable for magnets for compactmotors and sensors.

FIG. 4 is an isometric view of an embodiment using the hard magneticalloy in accordance with the present invention as a magnet for a hallpotentiometer. In FIG. 4, numeral 14 represents a magnetic sectioncomposed of the hard magnetic alloy in accordance with the presentinvention, and numeral 15 represents a supporting section for supportingthe magnetic section 14. The magnetic section 14 has a fan shape havinga permeance factor of approximately 5 and an absolute value of thetemperature coefficient of magnetization of 0.13%/° C. or less. Thesupporting section 15 includes a disk section 17 with a cutout section16 for placing the magnetic section 14, and a connecting column 18 whichprotrudes from the top face of the disk section 17.

In such a magnet for a hall potentiometer, the magnetic section 14 iscomposed of the hard magnetic alloy in accordance with the presentinvention, hence it has temperature-dependent properties which are equalto or superior to those of conventional ferrite magnets andNd-Fe-B-based magnets and the drift due to the change in temperature isprevented. As a result, the magnet enables accurate adjustment of thecircuit voltage in electronic devices. Further, the magnet in accordancewith this embodiment is more inexpensive than conventional ferritemagnets and Nd-Fe-B-based magnets, and has more excellent hard magneticcharacteristics than conventional ferrite magnets and Alnico magnets.

FIG. 5 is an isometric view of an embodiment using the hard magneticalloy in accordance with the present invention as a rotary encodermagnet 19. The rotary encoder magnet 19 is composed of the hard magneticalloy in accordance with the present invention and has a disk shape sothat the permeance factor is approximately 2. The rotary encoder magnet19 is magnetized along the circumference of the disk so as to formmultiple poles. Further, the rotary encoder magnet 19 has an absolutevalue of the temperature coefficient of magnetization of 0.15%/° C. orless.

The rotary encoder magnet 19 using the hard magnetic alloy in accordancewith the present invention has temperature-dependent properties whichare equal or superior to those of conventional ferrite magnets andNd-Fe-B-based magnets, can prevent the output drift due to a change intemperature, and thus can detect the rotated angle, or the like, of anelectronic device with high accuracy. Further, the rotary encoder magnet19 in accordance with this embodiment is more inexpensive thanconventional ferrite magnets and Nd-Fe-B-based magnets, and has moreexcellent hard magnetic characteristics than conventional ferritemagnets and Alnico magnets.

FIG. 6 is a cross-sectional view of a first embodiment using the hardmagnetic alloy in accordance with the present invention as a speakermagnet, wherein numeral 21 represents a pole piece composed of iron,numeral 22 is a cylindrical yoke which is provided at the exterior ofthe pole piece 21 at a distance, numerals 23 and 25 represent magnetscomposed of the hard magnetic alloy in accordance with the presentinvention and are provided at upper and lower portions of a gap betweenthe pole pieces 21 and the yoke 22, and numeral 25 represents avibrating cone. The magnets 23 and 24 have a ring shape. An audio coilnot shown in the drawing is provided between the magnetic gap formed bythese magnets 23 and 24, and is connected to the vibrating cone 25. Theaudio coil is vibrated by audio currents through an amplifier and causesthe vibrational movement of the vibrating cone 25, generating sound.

The magnets 23 and 24 composed of the hard magnetic alloy in accordancewith the present invention have temperature-dependent properties whichare equal or superior to those of conventional ferrite magnets andNd-Fe-B-based magnets, can prevent output drift due to a change intemperature, and thus can lead audio currents into the voice coil withhigh accuracy. Further, the speaker magnet in accordance with thisembodiment is more inexpensive than conventional ferrite magnets andNd-Fe-B-based magnets, and has more excellent hard magneticcharacteristics than conventional ferrite magnets and Alnico magnets.

FIG. 7 is a cross-sectional view of a second embodiment using the hardmagnetic alloy in accordance with the present invention as the magnetfor the speaker, wherein numerals 31 and 32 represent a pair of ironpole pieces which face in the vertical direction, numeral 33 representsa magnet composed of the hard magnetic alloy in accordance with thepresent invention and is provided between the pole pieces 31 and 32,numeral 34 represents a cylindrical yoke at the exterior of the polepieces 31 and 32 and the magnet 33 with an interval, numeral 35represents a vibrating cone, and numeral 36 represents amagnet-shielding cover. The magnet 33 has a ring shape. The pole pieces31 and 32 and the magnet 33 are fixed to the magnet-shielding cover 36with a bolt 37, washer 38 and nut 39. The speaker magnet 33 in thesecond embodiment composed of the hard magnetic alloy in accordance withthe present invention has substantially the same advantages as thespeaker magnet in the first embodiment.

Examples of preferred compositions of the hard magnetic alloy inaccordance with the present invention include Fe₈₈Pr₇B₅, Fe₈₆Nb₂Pr₇B₅,Fe₈₆Nd₇Zr₂B₅, Fe₈₆Nd₉B₅, Fe₈₄Pr₁₁B₅, Fe₈₈Nb₂Pr₅B₅, Fe₈₈Nb₂Nd₅B₅,Fe₈₆Nb₂Nd₇B₅, Fe₈₉Nb₂Pr₄B₅, Fe₈₉Nb₂Nd₄B₅, Fe₈₉Nb₂Pr₄B₅, Fe₉₀Nb₂Nd₅B₃,Fe₉₀Nb₂Pr₅B₃, Fe₈₉Nb₂Nd₅B₄, Fe₈₉Nb₂Pr₅B₄, Fe₈₄Nb₂Pr₇B₅Si₂,Fe₇₆Co₁₀Nb₂Pr₇B₅, and Fe₆₆Co₂₀Nb₂Pr₇B₅. These alloys can form a rigidcompact by annealing, and the formed fine crystalline phase contains abcc-Fe phase and an R₂Fe₁₄B phase. As a result, the alloys function aspermanent magnets having excellent hard magnetic characteristics.

EXAMPLES Example 1

Hard magnetic alloys were produced by annealing quenched alloy ribbonshaving various compositions.

Ingots were produced by an arc melting process, and the melted ingotswere sprayed onto a copper roll rotating in an Ar atmosphere through aslit nozzle with a slit width of approximately 0.3 mm to preparequenched alloy ribbons with a width of approximately 20 μm.

The resulting alloy ribbons were heated in an infrared image furnace of1×10⁻² Pa or less at a heating rate of 180° C./min. and annealed at anannealing temperature of 750° C. for 180 seconds to prepare annealedalloy ribbons of Example 1 in accordance with the present invention. Thecompositions of these alloy ribbons were Fe₇₆Co₁₀Nb₂Pr₇B₅,Fe₆₆Co₂₀Nb₂Pr₇B₅, and Fe₈₄Nb₂Pr₇B₅Si₂, respectively. FIGS. 8 to 10 aregraphs of demagnetization curves in the second quadrant of the resultingalloy ribbons which were determined with a vibrating sample magnetometer(VSM) in an applied magnetic field of 10 kOe in a vacuum at roomtemperature to approximately 217° C. In FIGS. 8 to 10, line A representsa permeance factor p of 10 corresponding to a prismatic shape, and lineB represents a permeance factor p of 1.5 corresponding to a disk shape.FIG. 11 is a graph illustrating changes in the remanent magnetization Irand the coercive force iHc with temperature of the alloy ribbons whichare determined by the demagnetization curves in the second quadrant.Table 1 shows magnetic characteristics of these alloy ribbons at roomtemperature. In Table 1, Ir/Is represents the ratio of the remanentmagnetization Ir to the remanence magnetization Is, that is a remanenceratio. Table 2 shows the temperature coefficients of Ir and iHc of eachalloy ribbon at a temperature of room temperature to 217° C., and thetemperature coefficients of Ir when the alloy has shapes causing p=1.5and p=10.

TABLE 1 Magnetic Characteristics (Room Temperature) Composition of hardmagnetic Is Ir material (emu · g⁻¹) (emu · g⁻¹) iHc (kOe) Ir/IsFe₇₆Co₁₀Nb₂Pr₇B₅ 147.2 114.8 3.1 0.78 Fe₆₆Co₂₀Nb₂Pr₇B₅ 156.3 127.1 3.110.81 Fe₈₆Nb₂Pr₇B₅Si₂ 139.3 105.2 3.62 0.75

TABLE 2 Temperature Coefficients of Ir and iHc (Room temperature to Ca.217° C.) dI/ dI/ dIr/ diHc/ dT(%/ dT(%/ Composition of dT dT °C) °C)hard magnetic (%/ (%/ at P = at P = material °C) °C) 1.5 10 ExampleFe₈₈Nb₂Pr₅B₅ −0.06 −0.43 −0.38 −0.12 Fe₇₆Co₁₀Nb₂Pr₇B₅ −0.02 −0.28 −0.20−0.10 Fe₆₆Co₂₀Nb₂Pr₇B₅ −0.02 −0.36 −0.33 −0.08 Fe₈₄Nb₂Pr₇B₅Si₂ −0.02−0.32 −0.17 −0.05 Comparative Fe₇₇Nd₁₅B₈ −0.16 −0.4 — — Example(Fe_(0.9)Co_(0.1))₇₇Nd₁₅B₈ −0.11 −0.35 — — Ferrite magnet −0.18 — — —

Relationships between the magnetic characteristics and the temperatureof a conventional ferrite magnet and a conventional Nd-Fe-B-based(Nd₂Fe₁₄B) magnet are also shown as Comparative Examples. Thetemperature coefficients of Ir and iHc of these conventional magnets arealso shown in Table 2.

The results shown in FIG. 11 and Table 2 demonstrate that the alloyribbons in accordance with the present invention and the magnets of thecomparative examples have a trend towards decreasing Ir and iHc withincreased temperature. The temperature coefficient of iHc (diHc/dt) ofthe alloy ribbon of the Example having a composition of Fe₈₈Nb₂Pr₅B₅ is−0.43%/° C. which is very near that of the Nd-Fe-B-based magnet of theComparative Example, that is, 0.4%/° C. The temperature coefficients ofthe alloy ribbons of the Example containing Co or Si, that is −0.28 to−0.36%/° C. are smaller than that of the Nd-Fe-B-based magnet. Such adecrease in the temperature coefficient of iHc by the addition of Co isprobably due to an increase in the Curie temperature of the hardmagnetic phase.

The temperature coefficient of Ir (dIr/dT) of the alloy ribbon of theExample having a composition of Fe₈₈Nb₂Pr₅B₅ is −0.06%/° C., which islow er than −0.11 to −0.16%/° C. of the Nd-Fe-B-based magnets ofComparative Examples having compositions of Fe₇₇Nd₅B₅ and(Fe_(0.9)Co_(0.1))₇₇Nd₁₅B₈, respectively. These results are probably dueto the following reasons; the alloy ribbon of the Example have anano-composite texture composed of a hard magnetic phase and a softmagnetic (bcc-Fe) phase having a small temperature coefficient ofmagnetization, whereas the magnets of the Comparative Examples have onlythe hard magnetic phase. The alloy ribbons containing Co or Si of theExample have a significantly low temperature coefficient of Ir, −0.02%/°C. FIG. 9 shows that the alloy ribbon having a composition ofFe₆₆Co₂₀Nb₂Pr₇B₅ has a particularly small change in magnetization withtemperature at a region of p=10 or more.

FIG. 12 is a graph of Ir of the alloy ribbons of the Example having ashape causing p=1.5 or p=10 at various temperatures, wherein the Irvalue is determined from the demagnetization curves shown in FIGS. 8 to10. FIG. 12 also shows the Ir values of a conventional Sm-Co-basedmagnet and a conventional Nd-Fe-B-based magnet having a composition ofNd₂Fe₁₄B for comparison, wherein these magnets are used in a shapecausing P=1.5 or P=10.

FIG. 13 is a graph illustrating the relationship between the permeancefactor and the temperature coefficient of the sintered bulk or alloycompact having a composition of Fe₈₈Nb₂Nd₅B₅ and an alloy ribbon havinga composition of Fe₈₆Nb₂Pr₇B₅ in accordance with the present invention,as well as a conventional Nd-Fe-B-based magnet for comparison having acomposition of Nd₂Fe₁₄B.

Table 2 and FIGS. 11 to 13 demonstrate that, at a low permeance factor,that is, p=1.5, the temperature coefficient of the Fe₈₄Nb₂Pr₇B₅Si₂ alloyribbon containing Si is 0.17%/° C., the temperature coefficients of theFe₇₆Co₁₀Nb₂Pr₇B₅ and Fe₆₆Co₂₀Nb₂Pr₇B₅ alloy ribbons containing Co, are−0.20%/° C. and −0.33%/° C., respectively, and the temperaturecoefficient of the Fe₈₈Nb₂Pr₅B₅ alloy ribbon is −0.38%/° C. Accordingly,alloy ribbons of the Example have relatively high temperaturecoefficients which are similar to that of the Comparative Example at thelow permeance. At a high permeance of p=10, the temperature coefficientof the Fe₈₈Nb₂Nd₅B₅ alloy is −0.12%/° C., and is similar to that of theconventional Nd-Fe-B-based magnet. The Fe₈₄Nb₂Pr₇B₅Si₂ alloy containingSi and the Fe₆₆Co₂₀Nb₂Pr₇B₅ alloy containing Co have smaller changes inmagnetic characteristics, that is, temperature coefficients of −0.05%/°C. and −0.08%/° C., respectively. The alloy ribbons of the Example, andparticularly the Fe₆₆Co₂₀Nb₂Pr₇B₅ alloy ribbon has a small absolutevalue of the temperature coefficient within a practical temperaturerange of 27° C. to 157° C. when they are used as a shape causing p=10,which is similar to that of the Sm-Co-based magnet and smaller than thatof the Nd₂Fe₁₄B magnet of the Comparative Example.

Accordingly, the results shown in Table 2 and FIG. 11 to 13 demonstratethat each alloy of the Example has an absolute value of the temperaturecoefficient which is similar to or smaller than that of the Nd₂Fe₁₄Bmagnet of the Comparative Example when they are used at a permeancefactor of 2 or more, and is smaller than 0.1%/° C. and that of theNd₂Fe₁₄B magnet of the Comparative Example when they are used at apermeance factor of 10 or more.

Example 2

A quenched amorphous alloy ribbon having a composition of Fe₈₈Nb₂Pr₅B₅was annealed at an annealing temperature of 750° C. and the change inmagnetization with temperature of the hard magnetic alloy was measured.Also, a quenched amorphous alloy ribbon having a composition ofFe₈₈Pr₇B₅ was annealed at an annealing temperature of 650° C. and thechange in magnetization with temperature of the hard magnetic alloy wasmeasured. These results are shown in FIG. 14. As shown in FIG. 14, eachcurve of magnetization decreasing with temperature has two steps, thisdemonstrates that two phases participate in magnetization of the hardmagnetic alloy. The step near 315° C. indicates the Curie temperature ofthe Nd₂Fe₁₄B phase, and the step near 770° C. indicates the Curietemperature of the bcc-Fe phase. No magnetization step due to anamorphous phase is found in the graph in FIG. 14, probably due to lowmagnetization and low volume fraction of the amorphous phase.

The magnetization curves in the second quadrant of the Fe₈₈Nb₂Pr₅B₅ andFe₈₈Pr₇B₅ hard magnetic alloys are shown in FIG. 15. The magnetizationcurves in FIG. 15 have no steps, that are observed in a magneticmaterial composed of a single phase. Accordingly, an exchange-coupledmagnet, in which a fine soft or semi-hard magnetic phase and a hardmagnetic phase is magnetically coupled, is obtained.

Example 3

Hard magnetic alloys were produced at various heating rates in theannealing step of quenched alloy ribbons having various compositions.

Ingots were produced by an arc melting process, and the melted ingotswere sprayed onto a copper roll rotating in an Ar atmosphere through aslit nozzle with a slit width of approximately 0.3 mm to preparequenched alloy ribbons with a thickness of approximately 20 μm.

The resulting alloy ribbons were heated in an infrared image furnace of1×10⁻² Pa or less at a heating rate of 3 to 216° C./min. and annealed atan annealing temperature of 700° C. to 750° C. for approximately 180seconds to prepare annealed alloy ribbons of Example 3 to observe thedependence of magnetic characteristics and nano-crystalline structures.The compositions of these alloy ribbons were Fe₈₈Nb₂Pr₅B₅ andFe₈₈Nb₂Nd₅B₅, respectively. The results are shown in Tables 3 to 6,wherein Ir/Is represents the ratio of the remanent magnetization to thesaturation magnetization. The magnetic characteristics of the alloyribbons were determined with a vibrating sample magnetometer (VSM) in anapplied magnetic field of 15 kOe at room temperature. In some alloys,the magnetization is not saturated at an applied magnetic field of 15kOe, however, this is regarded as a saturated magnetization in thefollowing description. In analysis of the nano-crystalline structure,the lattice constant and average crystal grain size of the bcc-Fe phasewere calculated from the results of diffraction angles and half-widthvalues of the X-ray diffraction peaks using the Sheller's equation,respectively.

TABLE 3 Fe₈₈Nb₂Pr₅B₅ (700° C.) Maximum magnetic Average Heating ResidualCoercive energy Lattice crystal rate magnetization force productconstant grain size Sample No. (° C./min.) (T) Ir/Is kOe (kJ/m³) (nm)(nm) 1(Example) 180 1.36 0.77 2.17 85 0.2871 35 2(Example) 102 1.45 0.821.79 74 0.2872 36 3(Example) 36 1.38 0.78 2.33 82 0.28695 43 4(Example)18 1.39 0.78 1.90 67 0.28719 46 5(Comparative 6 1.33 0.76 1.99 590.28731 48 Example) 6(Comparative 3 1.28 0.73 1.88 55 — — Example)

TABLE 4 Fe₈₈Nb₂Pr₅B₅ (750° C.) Maximum magnetic Average Heating ResidualCoercive energy Lattice crystal rate magnetization force productconstant grain size Sample No. (° C./min.) (T) Ir/Is kOe (kJ/m³) (nm)(nm) 7(Example) 180 1.37 0.76 2.44 81 0.28727 35.2 8(Example) 102 1.280.74 1.91 52 0.28727 34.7 9(Example) 36 1.33 0.74 2.43 75 0.28707 37.210(Example) 18 1.26 0.70 1.64 37 0.28714 45.6 11(Comparative 6 1.15 0.641.32 23 0.28711 46 Example) 12(Comparative 3 1.11 0.62 1.26 21 — —Example)

TABLE 5 Fe₈₈Nb₂Nd₅B₅ (700° C.) Maximum magnetic Average Heating ResidualCoercive energy Lattice crystal rate magnetization force productconstant grain size Sample No. (° C./min.) (T) Ir/Is kOe (kJ/m³) (nm)(nm) 13(Example) 180 1.23 0.75 1.86 74 0.2871 27 14(Example) 102 1.340.81 1.90 81 0.28682 26 15(Example) 36 1.32 0.79 1.85 75 0.28705 3316(Example) 18 1.17 0.70 1.54 39 0.28719 42 17(Comparative 6 1.22 0.732.00 63 0.2875 37 Example) 18(Comparative 3 1.05 0.69 1.82 23 — —Example)

TABLE 6 Fe₈₈Nb₂Nd₅B₅ (750° C.) Maximum magnetic Average Heating ResidualCoercive energy Lattice crystal rate magnetization force productconstant grain size Sample No. (° C./min.) (T) Ir/Is kOe (kJ/m³) (nm)(nm) 19(Example) 180 1.5 0.81 2.46 119 0.2873 28.9 20(Example) 102 1.480.80 2.37 102 0.28714 32.4 21(Example) 36 1.32 0.71 2.05 54 0.28714 37.422(Example) 18 1.36 0.73 2.21 83 0.28715 29.5 23(Comparative 6 1.35 0.722.16 53 0.2873 35.1 Example) 24(Comparative 3 1.31 0.70 2.01 39 — —Example)

FIGS. 16 and 17 are graphs illustrating the dependence of the latticeconstant and average crystal grain size on temperature of the bcc-Fephase of each of the alloys having the compositions of Fe₈₈Nb₂Pr₅B₅ andFe₈₈Nb₂Nd₅B₅ after annealing at 700° C. to 750° C. FIGS. 18 and 19 aregraphs illustrating the dependence of the remanent magnetization Ir,remanence ratio Ir/Is, coercive force iHc, and maximum magnetic energyproduct (BH)_(max) on temperature of each of the alloys having thecompositions of Fe₈₈Nb₂Pr₅B₅ and Fe₈₈Nb₂Nd₅B₅ which were annealed at700° C. to 750° C. for 5 minutes and then cooled at a cooling rate whichwas the same as the heating rate. FIG. 20 is a graph illustrating therelationship between the magnetic characteristics and the averagecrystal grain sizes, determined from the half-width of the X-raydiffraction pattern, of the bcc-Fe phase of the alloy ribbons havingcompositions of Fe₈₈Nb₂(Pr,Nd)₅B₅. In FIGS. 16 to 20, Ta represents theannealing temperature, and each dotted line in FIGS. 16 and 17represents the relationship between the lattice constant of the bcc-Fe(pure Fe) and the heating rate, for comparison.

As shown in Tables 3 to 6 and FIGS. 16 and 17, the lattice constants ofthe bcc phases in the Fe₈₈Nb₂Pr₅B₅ and Fe₈₈Nb₂Nd₅B₅ alloy ribbonsannealed at 700° C. to 750° C. do not depend on the heating rate, andare greater than the lattice constant, 0.28665 nm of the pure Fe,whereas the average crystal grain sizes of the bcc-Fe phases decreaseswith the heating rate. The average crystal grain sizes of the bcc-Fephases in the Fe₈₈Nb₂Pr₅B₅ and Fe₈₈Nb₂Nd₅B₅ alloy ribbons annealed atthe maximum heating rate, 216° C./min., are approximately 35 nm, and 27to 29 nm, respectively.

As shown in Tables 3 and 4, and FIG. 18, both the Ir values and theIr/Is ratios of the Fe₈₈Nb₂Pr₅B₅ alloy ribbons annealed at 700° C. andat 750° C. increase with the heating rate during the annealing. Theincrease is particularly noticeable in the annealing at 750° C. comparedwith the annealing at 700° C. Although the iHc value of the Fe₈₈Nb₂Pr₅B₅alloy ribbon is not dependent on the heating rate when the ribbon isannealed at 700° C., but is dependent on the heating rate when annealedat 750° C. Such increases in Ir and iHc values are probably because ofimproved exchange coupling characteristics caused by the decreasedaverage crystal grain size. No distinct difference in the dependence ofthe average crystal grain size of the bcc-Fe phase on the heating ratebetween the alloy ribbon annealed at 700° C. and the alloy ribbonannealed at 750° C. is, however, observed. It is therefore presumed thatlarge dependence of Ir and iHc values on the heating temperature in thealloy ribbon annealed at 750° C. is due to any factors other than thecrystal grain size of the bcc-Fe phase. The dependence of the (BH)_(max)value on the heating rate is observed in the Fe₈₈Nb₂Pr₅B₅ alloy ribbonsannealed at 700° C. and 750° C., and noticeably in the alloy ribbonannealed at 750° C.

As shown in Tables 3 and 4 and FIG. 19, both the Ir values and Ir/Isratios of the Fe₈₈Nb₂Nd₅B₅ alloy ribbons annealed at 700° C. and 750° C.increase when the heating rate is increased during the annealing,although the increases are not noticeable compared with the Fe₈₈Nb₂Pr₅B₅alloy ribbons. The iHc value of this alloy ribbon substantially does notchange with the heating rate in these annealing temperatures. On theother hand, the (BH)_(max) values are dependent on the heating rate inthe Fe₈₈Nb₂Pr₅B₅ alloy ribbons annealed at 700° C. and more noticeablyat 750° C. In the Fe₈₈Nb₂Pr₅B₅ alloy ribbons annealed at 700° C. and750° C., the crystal grain size of the bcc-Fe phase is small as a whole,hence the crystal grain size does not greatly contribute to the magneticcharacteristics, such as Ir and iHc values.

As shown in FIG. 20, the Ir/Is ratio and (BH)_(max) tend to decreasewith the increased crystal grain size of the bcc-Fe phase in both theFe₈₈Nb₂Pr₅B₅ and Fe₈₈Nb₂Nd₅B₅ alloy ribbons.

These experimental results suggest that the alloy ribbon in accordancewith the present invention heated at a high heating rate and annealed ata temperature of 700° C. to 750° C. has excellent hard magneticcharacteristics and a fine crystalline structure with an average crystalgrain size of 100 nm or less. The average crystal grain size of thebcc-Fe phase decreases at a heating rate of 10° C. or More and hardmagnetic characteristics such as iHc and (BH)_(max) are improved, hencethe heating rate during the annealing is set to 10° C./min. or more inthe present invention. The heating rate is preferably 36° C./min. ormore for achieving more excellent hard magnetic characteristics, andmore preferably 100° C./min. or more. In the Fe₈₈Nb₂Pr₅B₅ andFe₈₈Nb₂Nd₅B₅ alloy ribbons heated at a heating rate of 100° C./min. ormore, the crystal grain sizes of the bcc-Fe phase are approximately 35to 36 nm, and approximately 26 to 43.5 nm, respectively, and thecoercive forces are approximately 1.8 to 2.45 kOe, and approximately1.85 to 2.45 kOe, respectively.

Example 4

Quenched alloy ribbons were annealed at various annealing times toobtain hard magnetic alloy ribbons.

A quenched alloy ribbon with a thickness of approximately 20 μm wasprepared as in Example 1, and the quenched alloy ribbon was cut toprepare twenty alloy ribbon cut pieces with a thickness of 2 cm. The cutpieces were heated in an infrared image furnace under a pressure of1×10⁻⁵ Pa or less at a heating rate of 18° C./min. and annealed at anannealing temperature of 700° C. to 800° C. for 3 to 60 minutes toobserve change in magnetic characteristics. The alloy ribbon had acomposition of Fe₈₄Nb₂Nd₇B₅Si₂ in accordance with the present invention.The results are shown in Table 7. FIG. 21 is a graph illustrating thedependence of magnetic characteristics of the alloy ribbon on theannealing time. FIGS. 22 to 24 are graphs illustrating the dependence ofX-ray diffraction patterns of the alloy ribbons, which are annealed atvarious annealing temperatures, on the annealing time, wherein the X-raydiffraction patterns were measured with a Cu-Kα X-ray diffractometer.

TABLE 7 iHc (kOe) iHc (kOe) iHc (kOe) iHc (kOe) Annealing at a holdingat a holding at a holding at a holding temperature time of time of timeof time of (°C) 3 min. 10 min. 30 min. 60 min. 700 0.48 1.56 3.2 3.36750 3.44 3.5 3.09 2.89 800 2.93 2.78 2.6 2.51

As shown in Table 7 and FIG. 21, the maximum coercive force, 3.5 kOe, ofthe Fe₈₄Nb₂Nd₇B₅Si₂ alloy ribbon heated at a heating rate of 18° C./min.is obtained when the quenched alloy ribbon is annealed at 750° C. for 60min. Excellent hard magnetic characteristics, that is, a coercive forceof 3.38 kOe is achieved under annealing conditions of 700° C. for 60min. and 750° C. for 3 min. When the annealing temperature is 800° C.,the coercive force is not greater than 3 kOe at any of the annealingtimes.

The X-ray diffraction pattern, shown in FIG. 22, of the Fe₈₄Nb₂Nd₇B₅Si₂alloy ribbon heated at the heating rate of 18° C./min. and annealed at700° C. suggests that at a shorter annealing time of 3 to 30 minutes theFe₃B phase not contributing to hard magnetic characteristics isprecipitated, but the Nd₂Fe₁₄B phase having hard magneticcharacteristics is not substantially precipitated. The coercive force istherefore not high as shown in Table 7 and FIG. 21. On the other hand,at a long annealing time of 60 minutes, the Nd₂Fe₁₄B phase isprecipitated, and a high coercive force is achieved.

The X-ray diffraction pattern, shown in FIG. 23, of the Fe₈₄Nb₂Nd₇B₅Si₂alloy ribbon heated at the heating rate of 18° C./min. and annealed at750° C. suggests that at all annealing times the Nd₂Fe₁₄B phase isprecipitated. In particular, at annealing times of 3 minutes and 10minutes, high hard magnetic characteristics are obtained as shown inTable 7 and FIG. 21, whereas at longer annealing times of 30 minutes and60 minutes, the hard magnetic characteristics decrease, because thecrystal grains of the soft magnetic phase (bcc-Fe phase) continuegrowing during the excessive annealing time over the optimum annealingtime of 3 to 10 minutes.

The X-ray diffraction pattern, shown in FIG. 24, of the Fe₈₄Nb₂Nd₇B₅Si₂alloy ribbon heated at the heating rate of 18° C./min. and annealed at800° C. suggests that at all annealing times the Nd₂Fe₁₄B phase isprecipitated. The hard magnetic characteristics are, however, notexcellent as shown in Table 7 and FIG. 21, because of the crystal graingrowth in the soft magnetic phase (bcc-Fe phase) due to such a highannealing temperature.

Example 5

Hard magnetic alloy ribbons were produced by changing annealingconditions.

A quenched alloy ribbon with a thickness of approximately 20 μm having acomposition of Fe₈₄Nb₂Nd₇B₅Si₂ was prepared as in Example 1, and twentycut pieces with a length of 2 cm (the total weight of approximately 20mg) were produced from the quenched alloy ribbon. The twenty cut piecesand four quenched alloy ribbons (the total weight of 10 g) were annealedin an infrared image furnace with a pressure of 1×10⁻⁵ Pa or less underthe following annealing conditions for achieving satisfactory hardmagnetic characteristics, (1) at 700° C. for 60 minutes, (2) at 750° C.for 3 minutes, and (3) 750° C. for 10 minutes, wherein the heating ratewas 18° C./min. in these annealing conditions. The coercive forces ofthe annealed alloy ribbons are shown in Table 8.

TABLE 8 iHc(kOe) iHc(kOe) iHc(kOe) Size and number at 700° C., at 750°C., at 750° C., of alloy ribbon 60 min. 3 min. 10 min. 2 cm × 20 3.363.44 3.5 2.5 g × 4 3.73 3.33 3.49

Table 8 demonstrates that the Fe₈₄Nb₂Nd₇B₅Si₂ alloy ribbons annealedunder the above-mentioned conditions have high coercive forcesregardless of the total weight of the annealed ribbons. Accordingly, thepreferable annealing temperature ranges from 700° C. to 750° C., and thepreferable annealing time ranges from 3 minutes to 60 minutes.

Example 6

A quenched alloy ribbon having a composition of Fe₈₈Nb₂Nd₅B₅ inaccordance with the present invention, produced as in Example 3, washeated in an infrared image furnace of 1.3×10⁻³ Torr or less at aheating rate (α) of 3 to 180° C./min. and annealed at 700° C. to 750° C.for approximately 180 seconds to study the effect of the annealing onthe magnetic characteristics and nano-crystalline structure. Themicrostructure of the annealed alloy ribbon was observed with an X-raydiffractometer and a transmission electron microscope (TEM), andmagnetic characteristics were measured by a VSM at room temperaturewhile applying a magnetic field of 1.5 T to the alloy ribbon in thelongitudinal direction.

FIG. 25 is a graph illustrating the dependence of the demagnetizationcurves in the second quadrant of the Fe₈₈Nb₂Nd₅B₅ alloy ribbon, which isannealed at 750° C. for 180 seconds and quenched in the infrared imagefurnace, on the heating rate. FIG. 25 demonstrates that the heatingrates in accordance with the present invention, that is, 180° C./min.and 18° C./min. (Example) cause increased Ir, iHc and thus increased(BH)_(max) compared with the heating rate of 3° C./min. (ComparativeExample).

FIG. 26 is a graph illustrating the dependence of Ir, Ir/Is, iHc, and(BH)max of the resulting Fe₈₈Nb₂Nd₅B₅ alloy ribbon on the heatingtemperature. FIG. 26 demonstrates that the Ir, Ir/Is, iHc, and (BH)maxvalues of the alloy ribbon annealed at 750° C. for 180 secondssignificantly increase with the heating rate during the annealing.

FIG. 27 is a graph illustrating the dependence of Ir, Ir/Is, and iHcvalues of the Fe₈₈Nb₂Nd₅B₅ alloy ribbon on the heating temperature,which is heated at a heating rate of 3 to 180° C., annealed for 180seconds and quenched in the infrared image furnace. FIG. 27 demonstratesthat superior magnetic characteristics are achieved at an annealing rateof 750° C. In particular, the alloy in accordance with the presentinvention heated at a heating rate of 180° C./min. and annealed at atemperature of 700° C. to 800° C. has superior magnetic characteristicsto the alloy ribbon for comparison heated at a heating rate of 3°C./min.

FIG. 28 is a graph illustrating the dependence of magnetization ontemperature of the Fe₈₈Nb₂Nd₅B₅ alloy ribbon, which is heated at aheating rate of 3 to 180° C., annealed at a temperature of 750° C. for180 seconds, and quenched. FIG. 28 demonstrates that magnetization ofthe alloy ribbon decreases with the temperature rise regardless of theheating rate.

FIG. 29 is a graph illustrating the dependence of the average crystalgrain sizes of the bcc-Fe phase and the Nd₂Fe₁₄B phase on the heatingtemperature in the Fe₈₈Nb₂Nd₅B₅ alloy ribbon, which is annealed at 750°C. for 180 seconds and quenched, wherein the bcc-Fe (TEM) and theNd₂Fe₁₄B (TEM) refers to particle sizes determined from TEM photographs,and the bcc-Fe (XRD) refers the particle size determined from the X-raydiffraction pattern. FIG. 29 clearly shows that the average crystalgrain sizes of the bcc-Fe phase and the Nd₂Fe₁₄B phase decrease with theincreased heating rate. In particular, the Nd₂Fe₁₄B phase has a smalleraverage crystal grain size than that of the bcc-Fe phase, and is morelargely dependent on the heating rate than the bcc-Fe phase.

FIGS. 28 and 29 suggest that the dependence of the magneticcharacteristics of the Fe₈₈Nb₂Nd₅B₅ alloy ribbon on the heatingtemperature is originated from fining of the crystal grain size. Asmaller average grain size of the Nd₂Fe₁₄B crystalline phase causes anincreased probability of a bcc-Fe crystal grain and a Nd₂Fe₁₄B crystalgrain lying side by side, that is increased coupling of the bcc-Fecrystal grains with the Nd₂Fe₁₄B crystal grains. As a result, exchangecoupling characteristics are improved by exchange coupling between thebcc-Fe soft magnetic phase and the Nd₂Fe₁₄B hard magnetic phase.

FIG. 30 is a graph of DSC thermograms of an amorphous Fe₈₈Nb₂Nd₅B₅ alloyribbon obtained by a liquid quenching process at different heatingrates. The start point of the first peak in each DSC thermogramcorresponds to the crystallization temperature (T_(x1)) of the bcc-Fephase, and the start point of the second peak corresponds to thecrystallization temperature (T_(x2)) of the Nd₂Fe₁₄B phase. FIG. 31A isa graph illustrating the T_(x1) and T_(x2) determined from the DSCthermograms as shown in FIG. 30 vs. the heating rate, and FIG. 31B is agraph showing the difference between the T_(x1) and T_(x2),DT_(x)=T_(x2)−T_(x1), in FIG. 31A.

FIGS. 29 to 31 suggest that the difference DT_(x) decreases with theincreased heating rate during annealing, hence the average crystal grainsizes of both the bcc-Fe and Nd₂Fe₁₄B phases decrease with the increasedheating rate. At a higher heating rate, both the bcc-Fe and Nd₂Fe₁₄Bphases are simultaneously crystallized and prevent coarsening of otherphases.

Example 7

Hard magnetic alloy compacts were prepared as follows.

(Preparation of Amorphous Alloy)

Amorphous alloys having the compositions shown in Table 9 were preparedas follows. Alloy ingots based on the compositions were produced by anarc melting process, and each alloy melt was sprayed onto a copper rollrotating in an Ar atmosphere to prepare a quenched alloy ribbon with athickness of approximately 20 μm. The quenched alloy ribbon waspulverized with a rotor speed mill to prepare an amorphous alloy powderwith a particle size of 50 μm to 150 μm.

(Production of Alloy Compact)

The crystallization temperature T_(x) (° C.) of each amorphous alloypowder was determined by differential scanning calorimetry (DSC). Thealloy powder was placed into a mold and sintered with a hot press undera sintering pressure of 636 Mpa, and at a sintering time of 8 minuteswhile varying the sintering temperature T_(s) (° C.) to form an alloycompact.

(Measurement of Magnetic Characteristics)

The relative density (%) and the remanent magnetization Ir (T),remanence ratio Ir/Is, and coercive force iHc (kOe) of the resultingcompact were measured. The relative density (%) corresponds to the truedensity of approximately 7.5 g/cm³. The remanent magnetization Ir (T) isrepresented by the following equation:

Ir (T)=4π×7.5×(relative density)×Ir (emu/g)/10000

Table 9 shows the composition of each amorphous alloy, the sinteringtemperature T_(s) (° C.), the sintering pressure P_(s) (MPa), theannealing temperature (° C.), and saturated magnetization Is (T) values,remanent magnetization Ir (T) values, remanence ratios Ir/Is, coerciveforces iHc (kOe), maximum energy products (BH)_(max) and densities(g/cm³) in the X, Y, and Z direction in FIG. 32 when the sinteringpressure is applied in the Z direction.

TABLE 9 Alloy composition and sintering Annealing Is Ir iHc (BH)_(max)Density condition temperature Direction (T) (T) (kOe) (kJm⁻³) Ir/Is(g/cm³) Fe₉₀Nb₂Nd₅B₃ 750° C. X 1.598 1.146 2.862 64.8 0.717 7.31 Ts =600° C. Y 1.592 1.139 2.875 64.5 0.715 7.31 Ps = 636 MPa Z 1.573 1.1472.875 73.1 0.729 7.31 Fe₈₉Nb₂Nd₅B₄ 750° C. X 1.593 1.191 3.063 79.60.748 7.44 Ts = 600° C. Y 1.584 1.183 3.063 78.5 0.745 7.44 Ps = 636 MPaZ 1.583 1.17 3.063 79.8 0.739 7.44 Fe₈₉Nb₂Nd₄B₅ 750° C. X 1.636 1.2112.95 74.5 0.74 7.44 Ts = 600° C. Y 1.632 1.207 2.94 74.5 0.74 7.44 Ps =636 MPa Z 1.624 1.212 2.85 82 0.75 7.44 Fe₈₈Nb₂Nd₅B₅ 750° C. X 1.5211.109 3.525 78.6 0.729 7.70 Ts = 600° C. Y 1.52 1.109 3.512 78.8 0.737.70 Ps = 636 MPa Z 1.521 1.102 3.5 80.3 0.725 7.70 Fe₇₆Co₁₀Nb₂Nd₇B₅750° C. X 1.391 0.884 4.138 61.1 0.636 7.65 Ts = 600° C. Y 1.388 0.8864.125 61.6 0.638 7.65 Ps = 636 MPa Z 1.403 0.92 4.163 71.5 0.656 7.65Fe₈₄Nb₂Nd₇B₅Si₂ 750° C. X 1.284 0.83 4.813 55.7 0.646 7.55 Ts = 600° C.Y 1.286 0.831 3.55 55.6 0.646 7.55 Ps = 636 MPa Z 1.29 0.866 4.88 65.60.671 7.55 Fe₆₆Co₂Cr₅Nd_(5.5)B_(18.5) 700° C. X 1.176 0.797 6.338 63.40.678 7.51 Ts = 600° C. Y 1.175 0.797 6.325 64.1 0.678 7.51 Ps = 636 MPaZ 1.166 0.796 6.338 67 0.683 7.51 Ts = sintering temperature Ps =sintering pressure

The results shown in Table 9 demonstrate that dense alloy compactshaving excellent hard magnetic characteristics are produced bycompacting the amorphous alloy powders while crystallizing the amorphousalloys or growing the crystal grains. The Is, Ir, Ir/Is values arerelatively higher in the Z direction, and all the alloy compacts showhigh (BH)_(max) values. Accordingly, crystallization or crystal growthunder stress causes an anisotropic hard magnetic phase, and thusimproved hard magnetic characteristics.

FIG. 33A is a graph of the thermal mechanical analysis (TMA) thermogramof the Fe₈₈Nb₂Nd₅B₅ amorphous alloy ribbon, and FIG. 33B is a graph ofthe DSC thermogram of the amorphous alloy ribbon, wherein boththermograms were obtained at a heating rate of 0.33° C./sec. In the DSCthermogram in FIG. 33B, an exothermic peak observed at approximately850° C. suggests the crystallization of the bcc-Fe or FeB phase. On theother hand, as shown in the TMA thermogram in FIG. 33B, the elongationof the amorphous alloy ribbon significantly increases from thetemperature region near 427° C. which is approximately 200° C. lowerthan the crystallization temperature. Such a phenomenon suggests asoftening of the alloy near the crystallization temperature. FIG. 33Aalso includes a TMA thermogram of the crystalline alloy having the samecomposition for comparison. FIG. 33A suggests that softening of thecrystalline alloy does not occur.

Particles of the amorphous alloy powder are closely bonded to each otherby being pressurized at the softening temperature region to form a densealloy compact with decreased pores.

FIGS. 34A to 34C are schematic photographs observed by microscopy of thetexture in the alloy compact which is sintered at a sintering pressureof 636 MPa, a sintering time of 8 minutes, and sintering temperatures of400° C., 500° C. and 600° C., respectively. A more dense alloy compactwith decreased pores can be obtained when the sintering temperature israised, and an alloy compact sintered at 600° C. is sufficiently dense.

At least part of the amorphous phase is crystallized when being heatedto the crystallization temperature.

FIG. 35 is a graph of X-ray diffraction patterns of the alloy compactimmediately after sintering at a sintering pressure of 636 MPa, asintering time of 8 minutes, and various sintering temperatures. In FIG.35, patterns (a), (b), (c), and (d) correspond to sintering temperaturesof 400° C., 500° C., 600° C., and 650° C., respectively. In thesepatterns, the hallow pattern at 2θ=ca. 44.5° suggests the presence ofthe bcc-Fe crystalline phase.

FIG. 35 demonstrates that at a sintering temperature of 500° C. or lessonly the hallow pattern is observed as shown in patterns (a) and (b) andthus the bcc-Fe crystalline phase does not substantially form, whereasat a sintering temperatures of 600° C. to 650° C., a bcc-Fe crystallinephase, a Nd₂F₁₄B phase, and an Fe-B compound phase, supporting hardmagnetic characteristics, form as shown in patterns (c) and (d).

FIG. 36 is a graph of the density of the alloy compact which is sinteredat a sintering pressure of 636 MPa, a sintering time of 8 minutes, anddifferent sintering temperature. FIG. 36 demonstrates that the densityincreases with the sintering temperature, and a high density of 7.45g/cm³ or more is achieved at a sintering temperature of higher than 500°C.

The results shown in FIG. 36 in view of the X-ray diffraction patternssuggest that the amorphous alloy compact sintered at 500° C. or less hasa relatively low density, whereas the alloy compact sintered at 600° C.or more contains the bcc-Fe crystalline phase and has a sufficientlyhigh density. These results suggest that when the amorphous alloy powderis molded under stress by utilizing the softening phenomenon before thecrystallization, the alloy powder is simultaneously crystallized andsintered, and thus the alloy compact has a high density, and excellentmechanical and hard magnetic characteristics.

For comparison, a crystalline powder, obtained by crystallizing anamorphous powder at 750° C. and then annealing for 3 minutes, wassintered at a sintering pressure of 636 MPa, a sintering time of 8 min.and a sintering temperature of approximately 600° C. The density of thisalloy compact is also shown in FIG. 36. The alloy compact from thecrystalline powder has a small density, i.e., approximately 6.70 g/cm³.Accordingly, a dense alloy compact cannot be obtained by the sinteringof the crystalline powder.

FIG. 37 is a graph illustrating the density of the compact which isobtained at a sintering temperature of 600° C. or 650° C., and asintering pressure of 260 to 636 MPa, as well as the density of an alloycompact, which is produced by sintering a crystalline powder at asintering temperature of 600° C., at a sintering pressure of 636 MPa,and a holding time of 3 minutes, in which the crystalline powder isprepared by crystallizing an amorphous powder at 750° C. and thenannealing for 3 minutes.

FIG. 37 demonstrates that the density of the compact from the amorphousalloy powder increases with the sintering pressure and reaches the truedensity at a sintering pressure of 500 MPa or more. In contrast, thecompact from the crystalline powder has a significantly low densityregardless of a high sintering pressure of 636 MPa.

As a preferable embodiment of a method for producing the compact inaccordance with the present invention, an Fe-Nd-Nb-B-based alloy compactwas produced by a spark plasma sintering process using a sintering dicewith a punch diameter of 18 mm, and the magnetic characteristics of theresulting compact were observed.

(Preparation of Quenched Amorphous Alloy Ribbon)

Two amorphous alloys having compositions of Fe₈₈Nb₂Nd₅B₅ andFe₈₆Nb₂Nd₇B₅, respectively, were prepared as follows. An ingot of eachalloy was produced by an arc melting process, the alloy melt was sprayedonto a copper single roll rotating in an Ar atmosphere through a quartznozzle with a slit having a thickness of 0.3 mm and a width of 14 mm toobtain a quenched alloy ribbon having a thickness of approximately 20μm. The conditions for quenching the alloy melt were as follows:

Weight of the melt: 15 to 20 g

Degree of attained vacuum: 6×10³ Pa or less

Pressure of Ar atmosphere: 15 cmHg

Spraying pressure: 0.4 kgf/cm³

Rotation rate of roll: 4,000 rpm

Spraying temperature: 1,450° C.

(Preparation of Alloy Powder)

The quenched alloy ribbon was pulverized with a rotor speed mill, andthe alloy powder was classified. The particle size of the alloy powderprimarily ranges from 37 to 105 μm. In particular, the Fe₈₈Nb₂Nd₅B₅powder is primarily composed of 53 to 105 μm particles and theFe₈₆Nb₂Nd₇B₅ powder is primarily composed of 37 to 53 μm particles. Theresults suggest that the higher-Nd alloy is more fragile than thelower-Nd alloy, and can be easily pulverized.

For comparison, an Fe-(Nb,Zr)-B amorphous alloy ribbon containing norare earth metal was prepared and pulverized. The yield of the powderhaving a particle size ranging from 53 to 105 μm was less than 10%.

Accordingly, the Fe-Nb-Nd-B amorphous alloy can be easily pulverizedcompared with the Fe-(Nb,Zr)-B amorphous alloy.

(X-ray Diffractometry of Alloy Powder)

FIGS. 38 and 39 are graphs of X-ray diffraction patterns of theFe₈₈Nb₂Nd₅B₅ and Fe₈₆Nb₂Nd₇B₅ powders, respectively. The hallow peaknear 2θ=50°, observed in all the diffraction patterns, suggestsformation of the amorphous phase in the powders. For comparison, theX-ray diffraction patterns of the powder from the Fe₈₄Nb₇B₉ amorphousalloy ribbon containing no rare earth element were also observed.Diffraction peaks based on a crystalline phase were observed in thepowder with a particle size ranging from 25 to 53 μm. The amorphouspowder may be crystallized during pulverizing, or a crystalline materialis mixed in the pulverizer.

Accordingly, the Fe-Nb-Nd-B amorphous alloy in accordance with thepresent invention can be easily pulverized without deteriorating theamorphous phase, and the powder has a desirable particle sizedistribution.

(Production of Alloy Compact)

Each of the amorphous Fe₈₈Nb₂Nd₅B₅ and Fe₈₆Nb₂Nd₇B₅ powders with aparticle size range of 37 to 105 μm was placed between the upper andlower punches 2 and 3 of the spark plasma sintering system shown in FIG.3, and the system was evacuated so that the atmosphere in the chamber 13was 3×10⁻³ Pa or less, while pressing the punches 2 and 3 from bothsides and applying electrical current pulses for heating the powder. Thepowder was sintered together with crystallization or crystal graingrowth and a compact was formed under the conditions of a sinteringpressure of 636 MPa, a heating rate of 1.8° C./sec., a sinteringtemperature of 500° C. to 600° C., and a holding time of 8 minutes. Theresulting compact was annealed in an atmosphere of 1.3×10⁻³ Pa or lessat a heating rate of 3° C./sec., an annealing temperature of 550° C. to800° C., and an annealing time of 3 minutes.

FIG. 40 is a graph illustrating the relationship between the time, thetemperature of the powder, and the distance X between the two dices (orexpansion), wherein powders A and B represent the amorphous Fe₈₈Nb₂Nd₅B₅and Fe₈₆Nb₂Nd₇B₅ powders, respectively, powder C represents anano-crystallized Fe₈₈Nb₂Nd₅B₅ powder annealed at 750° C., and thesepowders were sintered at a heating rate of 1.8° C./sec. and a sinteringtemperature of 600° C. The temperature of the powder was measured with athermocouple 7 attached to the dice, and the distance X between thedices was determined by the distance between the upper and lower punches2 and 3, as shown in FIG. 1.

As shown in FIG. 40, the distance X of the nano-crystallizedFe₈₈Nb₂Nd₅B₅ powder (powder C) causes monotone increase with a rise intemperature, whereas the expansion of the amorphous Fe₈₈Nb₂Nd₅B₅ andFe₈₆Nb₂Nd₇B₅ powders (powders A and B) is completed at approximately327° C. or 240 seconds, and on the contrary, powders shrink at furthertimes. Such a phenomenon suggests that these amorphous alloy powders(powders A and B) are softened near their crystallization temperature,i.e., 327° C., respectively, resulting in the increased densities of thecompacts.

(Structure and Magnetic Characteristics of Compact)

FIG. 41 is a graph illustrating the densities of the compacts which wereformed by sintering the amorphous Fe₈₈Nb₂Nd₅B₅ and Fe₈₆Nb₂Nd₇B₅ powders,as well as the crystalline Fe₈₈Nb₂Nd₅B₅ powder for comparison, at asintering pressure of 636 MPa, a heating rate of 1.8° C./sec., asintering temperature of 500° C. to 600° C., and a holding time of 8minutes. As shown in FIG. 41, the density of the compacts obtained bysintering the amorphous Fe₈₈Nb₂Nd₅B₅ and Fe₈₆Nb₂Nd₇B₅ powders at arelatively high temperature of 600° C. is approximately 7.5×10⁻³ kg/m³,which is almost the same as the true density, whereas the density of thecompact obtained by the crystalline Fe₈₈Nb₂Nd₅B₅ powder at the samesintering temperature is approximately 6.6×10⁻³ kg/m³. Such highdensities of the amorphous alloy powders also suggests thesolidification of the amorphous alloy powder accompanied by softeningnear the crystallization temperature. The density of the Fe₈₆Nb₂Nd₇B₅compact slightly decreases at a lower sintering temperature, whereas theFe₈₈Nb₂Nd₅B₅ compact has a high density of 7.5×10⁻³ kg/m³ even at a lowsintering temperature of 500° C. According to DSC with a heating rate of0.67° C./sec., the Fe₈₈Nb₂Nd₅B₅ and Fe₈₆Nb₂Nd₇B₅ alloys havecrystallization temperatures of 619° C. and 643° C., respectively. Thereason for the Fe₈₈Nb₂Nd₅B₅ compact having a high density at arelatively low temperature is probably its low crystallizationtemperature.

FIG. 42 is a graph illustrating X-ray diffraction patterns of theFe₈₈Nb₂Nd₅B₅ and Fe₈₆Nb₂Nd₇B₅ compacts immediately after the sparkplasma sintering process, wherein each phase in the alloys wereidentified by the characteristic peaks of the X-ray diffraction patternsobtained by the Co-Kα line. As shown in FIG. 42, the diffraction patternof the Fe₈₈Nb₂Nd₅B₅ compact after sintering includes a peak (◯) based onthe bcc-Fe phase, peaks (□) based on the Fe₂B phase, and peaks (▪) basedon the Nd₂Fe₁₄B phase, and the diffraction pattern of the Fe₈₆Nb₂Nd₇B₅compact include peaks (Δ) based on the Fe₃B phase, as well as the peaksbased on the bcc-Fe phase, Fe₂B phase, and Nd₂Fe₁₄B phase. Thediffraction intensities of the peaks based on the crystalline phases ofeach amorphous compact increase with the sintering temperature, and thissuggests prompt crystallization. The Fe₈₆Nb₂Nd₇B₅ compact sintered at500° C. or 550° C. has a relatively broad diffraction pattern, whichsuggests a mixture of crystalline amorphous phases. Such a compact has arelatively low density, that is, 6.6 to 7.0×10⁻³ kg/m³ as shown in FIG.13. FIGS. 41 and 42 demonstrate that a high density of the compact isachieved by simultaneous solidification and crystallization of theamorphous alloy powder.

FIG. 43 is a graph illustrating the magnetic characteristics of theamorphous Fe₈₈Nb₂Nd₅B₅ and Fe₈₆Nb₂Nd₇B₅ compacts which were sintered ata sintering pressure of 736 Mpa, a heating rate of 1.8° C./sec., asintering temperature of 500° C. to 600° C., and a holding time of 8minutes, and are annealed at a heating rate of 3° C./min, at anannealing temperature of 750° C., and an annealing time of 180 seconds.The magnetic characteristics were measured with a VSM at roomtemperature while applying a magnetic field of 1.5 T to the test piece(a thickness of 0.5 mm, a width of 1.5 mm, and a length of 6 mm) in thelongitudinal direction. As shown in FIG. 43, the remanent magnetizationand remanence ratio of the Fe₈₈Nb₂Nd₅B₅ compact slightly changes withthe sintering temperature, whereas the remanent magnetization, remanenceratio, and coercive force of the Fe₈₆Nb₂Nd₇B₅ compact increase with thesintering temperature, and the compact has excellent magneticcharacteristics at a sintering temperature of 600° C.

FIG. 44 is a graph illustrating the relationship between the magneticcharacteristics and the sintering temperature of the Fe₈₈Nb₂Nd₅B₅compact which was sintered with a sintering dice having a diameter of 10mm at a sintering pressure of 636 Mpa. FIG. 44 demonstrates that thesaturation magnetization Is, remanent magnetization Ir, and theremanence ratio Ir/Is of the compact decrease with a decreased sinteringtemperature.

FIG. 45 is a graph illustrating the relationship between the magneticcharacteristics and the sintering pressure of the Fe₈₈Nb₂Nd₅B₅ compactwhich is sintered at a sintering temperature of 600° C. FIG. 45demonstrates that the saturation magnetization Is, remanentmagnetization Ir, and the remanence ratio Ir/Is of the compact decreasewith a decreased sintering pressure. As shown in FIGS. 36 and 37, thedensity of the compact decreases with a decreased sintering temperatureor decreased sintering pressure, hence a decreased density of thecompact will adversely affect the magnetic characteristics.

FIG. 46 is a graph illustrating the relationship between the magneticcharacteristics and the density of the Fe₈₈Nb₂Nd₅B₅ compact. As shown inFIG. 46, the saturation magnetization Is, remanent magnetization Ir, andremanence ratio Ir/Is decrease with a decreased density. Accordingly, ahigher density of the compact is essential for achieving higher magneticcharacteristics.

FIG. 47 is a graph of X-ray diffraction patterns of the Fe₈₈Nb₂Nd₅B₅ andFe₈₆Nb₂Nd₇B₅ compacts which were sintered at a sintering pressure of 736Mpa, a heating rate of 1.8° C./sec., and a holding time of 8 minutes,and were annealed at an annealing temperature of 750° C. As shown inFIG. 47, the diffraction pattern of the Fe₈₈Nb₂Nd₅B₅ compact aftersintering includes a peak (◯) based on the bcc-Fe phase, peaks () basedon the Nd₂Fe₁₄B phase, and peaks (Δ) based on the Fe₃B phase, and thediffraction pattern of the Fe₈₆Nb₂Nd₇B₅ compact include the peaks basedon the bcc-Fe phase and Nd₂Fe₁₄B phase. These compacts therefore containmixtures of soft and hard magnetic phases. The crystal grain sizes,determined from the half-width values of the (100) diffraction peak ofthe bcc-Fe phase of the Fe₈₈Nb₂Nd₅B₅ and Fe₈₆Nb₂Nd₇B₅ compacts areapproximately 20 nm, and 30 nm, respectively, which are the same asthose of the corresponding alloy ribbons. Also, in the high-resolutiontransmission electron microscope photographs of the Fe₈₈Nb₂Nd₅B₅ compactafter annealing at 750° C., bcc-Fe phases with grain sizes of 20 to 40nm and Nd₂Fe₁₄B phases with a grain size of 20 nm are observed. Theseresults illustrate the formation of the nano-crystalline compositetexture in the hard magnetic alloy compact in accordance with thepresent invention after annealing at 750° C., as in the alloy ribbon.

FIG. 48 is a graph illustrating magnetization curves of the Fe₈₈Nb₂Nd₅B₅(solid line) and Fe₈₆Nb₂Nd₇B₅ (broken line) compacts which were sinteredat a sintering pressure of 636 MPa, a heating rate of 1.8° C./sec., asintering temperature of 600° C., and a holding time of 8 minutes, andannealed at 750° C. These two compacts show smooth magnetization curvesas in a magnetic material composed of a single phase. Such amagnetization curve, which is shown in a hard magnetic material composedof a single hard magnetic phase, demonstrates that the fine soft orsemi-hard magnetic phase and the fine hard magnetic phase aremagnetically coupled with each other in the compact, and thus thecompact has exchange coupled magnetic characteristics.

The remanent magnetization Ir (T), remanence ratio Ir/Is, coercive forceiHc (kOe), and maximum energy product (BH)_(max) (kJ/m³) of theFe₈₈Nb₂Nd₅B₅ and Fe₈₆Nb₂Nd₇B₅ compacts were measured. For comparison,the magnetic characteristics of the alloy ribbons having the samecompositions and annealed at 750° C. were also measured. The results areshown in Table 10.

TABLE 10 Ir iHc (BH)_(max) Alloy Composition (T) Ir/Is (kOe) (kJm⁻³)Compact Fe₈₈Nb₂Nd₅B₅ 1.0 0.67 2.38 43 Fe₈₆Nb₂Nd₇B₅ 0.83 0.65 3.66 49Ribbon Fe₈₈Nb₂Nd₅B₅ 1.12 0.75 2.9 67 Fe₈₆Nb₂Nd₇B₅ 0.90 0.72 4.15 64

The results in Table 10 suggest that the alloy compacts have similarcoercive forces to those of the alloy ribbons having the samecompositions. The maximum magnetic energy products (BH)_(max) of thealloy compact are lower than those of the corresponding alloy ribbons,probably because of the decreased remanence ratios.

FIG. 49 is a graph illustrating the magnetization curves of an amorphousFe₈₈Nb₂Nd₅B₅ compact (solid line) and a crystalline Fe₈₈Nb₂Nd₅B₅ compact(broken line) which were sintered at a sintering pressure of 636 MPa, aheating rate of 1.8° C./sec., a sintering temperature of 600° C., and aholding time of 8 minutes, and annealed at a 750° C. As shown in FIG.49, these compacts have smooth magnetization curves with no steps, andthe amorphous Fe₈₈Nb₂Nd₅B₅ compact (solid line) has hard magneticcharacteristics superior to those of the crystalline Fe₈₈Nb₂Nd₅B₅compact (broken line).

FIGS. 50 and 51 are graphs of magnetic characteristics of the amorphousFe₈₈Nb₂Nd₅B₅ and Fe₈₆Nb₂Nd₇B₅ compacts which were annealed at asintering pressure of 636 MPa, a heating rate of 1.8° C./sec., asintering temperature of 500° C. to 600° C., and a holding time of 8minutes, and annealed at a heating rate of 3° C./min., at an annealingtemperature of 627° C. to 827° C., and an annealing time of 180 seconds.As shown in FIGS. 50 and 51, the Fe₈₈Nb₂Nd₅B₅ compact has a highcoercive force in a combination of a sintering temperature of 600° C.and an annealing temperature of 750° C., and the Fe₈₆Nb₂Nd₇B₅ compacthas a high coercive force in a combination of a sintering temperature of500° C. or 550° C. and an annealing temperature of 750° C. although thechange of the magnetic characteristics of the Fe₈₆Nb₂Nd₇B₅ compact isslight at a sintering temperature of 600° C.

In Table 10, the magnetic characteristics of the sintered compact areinferior to those of the alloy ribbon, because the demagnetization fieldis probably not corrected for the thickness of the sintered compact.

FIGS. 52 to 55 are graphs of demagnetization curves of the amorphouscompacts which were prepared by sintering the corresponding amorphousalloy powders at a sintering pressure of 636 MPa, a sinteringtemperature of 600° C., and a holding time of 8 minutes, and cutting thesintered alloys to form a cube of 5×5×5 mm or 4×4×4 mm, and annealingthe cube at the optimum annealing conditions. The demagnetization curveswere measured while applying a magnetic field of 5 T with a pulsemagnetization system. In each curve, the antimagnetic field iscorrected. FIG. 52 includes demagnetization curves in the X, Y, and Zdirections of the amorphous Fe₉₀Nb₂Nd₅B₃ compact, FIG. 53 includes thoseof the amorphous Fe₈₉Nb₂Nd₄B₅ compact, FIG. 54 includes those of theamorphous Fe₇₆Co₁₀Nb₂Nd₇B₅ compact, and FIG. 55 includes those of theamorphous Fe₈₄Nb₂Nd₇B₅Si₂ compact. The amorphous alloy powders werecompressed in the Z direction when they were sintered, hence theamorphous phase was crystallized or the crystal grains were grown underpressure. Each of these alloy compacts have a remanent magnetization of0.8 T or more, and a coercive force of approximately 2.5 kOe or more, ofwhich the hard magnetic characteristics are superior to those measuredwith a VSM in an applied magnetic field of 1.5 T. As shown in FIGS. 52to 55, each compact has a convex demagnetization curve in the Zdirection. Such excellent hard magnetic characteristics in the Zdirection are based on single-axis anisotropy of the hard magnetic phasewhich forms by the precipitation of the hard magnetic phase underpressure or stress.

Example 8

A quenched alloy ribbon having a composition of Fe₈₈Nb₂Nd₅B₅ wasprepared as in Example 1. The resulting alloy ribbon was heated to 750°C. in an infrared image furnace and annealed for approximately 180seconds. In the annealing process, as shown in FIG. 56, the heating rateof the alloy ribbon was 180° C./min. from room temperature (27° C.) toT₁, and 3° C./min. from T₁ to 750° C., wherein T₁ is changed within arange from 27° C. to 750° C. Several alloy ribbons annealed by differentheating processes were prepared in such a manner. Also, several annealedalloy ribbons were prepared by heating, as shown in FIG. 57, at aheating rate of 3° C./min. from room temperature (27° C.) to T₂, and180° C./min. from T₂ to 750° C., wherein T₂ is changed within a rangefrom 27° C. to 750° C.

FIG. 58 is a graph of DSC thermograms at a heating rate of 6 to 40°C./min. of the annealed alloy ribbons. In each of the heating rates, anexothermic peak based on the crystallization of the bcc-Fe phase wasobserved within a temperature range from 500° C. to 650° C. and anexothermic peak based on the crystallization of the Nd₂Fe₁₄B phase wasobserved in a temperature range above 650° C.

FIG. 59 is a graph illustrating the remanent magnetization Ir, remanenceratio Ir/Is, and coercive force iHc on T₁ of the annealed alloy ribbon,wherein T₁ was changed within a temperature range from 27° C. to 750° C.FIG. 59 demonstrates that all the remanent magnetization Ir, remanenceratio Ir/Is, and coercive force iHc are low at a T₁ of less than 500°C., resulting in deterioration of the hard magnetic characteristics.When T₁ is lower than 500° C., the heating rate is low in a temperaturerange (500° C. to 650° C.) causing the precipitation of the bcc-Fephase, hence the sites of the bcc-Fe nuclei will decrease and thecrystal grain size of the Nd₂Fe₁₄B phase will increase, which isprecipitated and grown between crystal grains of the bcc-Fe phase. Sucha phenomenon results in decreased exchange coupling, and thusdeterioration of hard magnetic characteristics.

When T₁ is 650° C., the heating rate is low in a temperature range (650°C. to 750° C.) causing the precipitation of the Nd₂Fe₁₄B phase. Althougha decrease in exchange coupling due to coarsening of the Nd₂Fe₁₄Bcrystal grains is expected, the observed hard magnetic characteristicsof the alloy ribbon are excellent as shown in FIG. 59. With such aphenomenon it is presumed that a large number of crystal grains of thebcc-Fe phase are precipitated at a high heating rate, hence the crystalgrain growth of the Nd₂Fe₁₄B phase is suppressed and thus exchangecoupling is improved.

FIG. 60 is a graph illustrating the remanent magnetization Ir, remanenceratio Ir/Is, and coercive force iHc on T₂ of the annealed alloy ribbon,wherein T₂ was changed within a temperature range from 27° C. to 750° C.FIG. 60 demonstrates that all the remanent magnetization Ir, remanenceratio Ir/Is, and coercive force iHc are low at a T₂ of less than 650°C., resulting in deterioration of hard magnetic characteristics. When T₂is lower than 650° C., the heating rate is low in a temperature range(500° C. to 650° C.) causing the precipitation of the bcc-Fe phase, thecrystal grain size of the bcc-Fe phase will increase, resulting indecreased exchange coupling with the Nd₂Fe14B phase and thusdeterioration of hard magnetic characteristics.

Accordingly, the alloy ribbon in accordance with the present inventionhas excellent hard magnetic characteristics by being annealed at aheating rate of 10° C./min. or more within a temperature range (500° C.t 650° C.) causing the precipitation of the bcc-Fe phase.

What is claimed is:
 1. A hard magnetic alloy compact comprising one ofan Fe-based and FeCo-based alloy containing 3 to 20 atomic percent of atleast one rare earth element R, and 2 to 20 atomic percent of B, whereinthe alloy having a texture, in which at least a part of the texturecomprises one of an amorphous phase and fine crystalline phase having anaverage crystal grain size of not greater than 100 nm, is subjected toone of crystallization and grain growth under stress, such that a mixedphase composed of a hard magnetic phase and one of a soft magnetic andsemi-hard magnetic phase is formed in said texture, anisotropy isimparted to a crystal axis of said hard magnetic phase, and said hardmagnetic alloy compact has a coercive force of not less than 1 kOe.
 2. Ahard magnetic alloy compact according to claim 1, wherein said hardmagnetic alloy compact comprises at least 10 percent by volume of one ofa soft magnetic and semi-hard magnetic phase having a coercive force ofnot greater than 1 kOe which comprises an Fe-B compound phase, anamorphous phase, and one of a body centered cubic (bcc) Fe phase andbcc-FeCo phase as precipitates, and at least 10 percent by volume of ahard magnetic phase having a coercive force of not less than 1 kOe whichcomprises an R₂Fe₁₄B phase.
 3. A hard magnetic alloy compact accordingto claim 1, wherein, after quenching, the alloy is compacted and issimultaneously one of crystallized and subjected to crystal growth understress.
 4. A hard magnetic alloy compact according to claim 1, whereinsaid hard magnetic alloy contains an amorphous phase, and is formed bysolidifying an alloy having hard magnetic characteristics when beingcrystallized, by means of a softening phenomenon of said alloy duringthe crystallization reaction.
 5. A hard magnetic alloy compact accordingto claim 1, wherein said alloy is heated under stress.
 6. A hardmagnetic alloy compact according to claim 1, wherein a relative densityof the compact obtained by compacting said alloy is not less than 90%.7. A hard magnetic alloy compact according to claim 1, wherein said hardmagnetic alloy compact has a remnant magnetization of not less than 100emu/g.
 8. A hard magnetic alloy compact according to claim 1, wherein aratio of a remnant magnetization Ir to a saturation magnetization Is isnot less than 0.6.
 9. A hard magnetic alloy compact according to claim1, wherein said hard magnetic alloy compact has the following formula:T_(x)M_(y)R_(z)B_(w) wherein T represents at least one element selectedfrom the group consisting of Fe, Co and Ni, M represents at least oneelement selected from the group consisting of Zr, Nb, Ta and Hf, Rrepresents at least one rare earth element, and the suffixes x, y, z andw by atomic percent satisfy 50≦x, 0≦y≦15, 3≦z≦20, and 2≦w≦20,respectively.
 10. A hard magnetic alloy compact according to claim 9,wherein the suffixes x, y, z and w by atomic percent satisfy 80≦x≦92,1≦y≦5, 3≦z≦10, and 3≦w≦7, respectively.
 11. A hard magnetic alloycompact according to claim 1, wherein said hard magnetic alloy compacthas the following formula: T_(x)M_(y)R_(z)B_(w)Si_(u) wherein Trepresents at least one element selected from the group consisting ofFe, Co and Ni, M represents at least one element selected from the groupconsisting of Zr, Nb, Ta and Hf, R represents at least one rare earthelement, and the suffixes x, y, z, w, and u by atomic percent satisfy50≦x, 0≦y≦15, 3≦z≦20, 2≦w≦20, and 0≦u≦5, respectively.
 12. A hardmagnetic alloy compact according to claim 11, wherein the suffixes x, y,z, w, and u by atomic percent satisfy 80≦x≦92, 1≦y≦5, 3≦z≦10, 3≦w≦7, and0.5≦u≦5, respectively.
 13. A hard magnetic alloy compact according toclaim 1, wherein said hard magnetic alloy compact has the followingformula: T_(x)M_(y)R_(z)B_(w)E_(v) wherein T represents at least oneelement selected from the group consisting of Fe, Co and Ni, Mrepresents at least one element selected from the group consisting ofZr, Nb, Ta and Hf, R represents at least one rare earth element, Erepresents at least one element selected from the group consisting ofCr, Al, Pt and platinum elements, and the suffixes x, y, z, w, and v byatomic percent satisfy 50≦x, 0≦y≦15, 3≦z≦20, 2≦w≦20, and 0≦v≦10,respectively.
 14. A hard magnetic alloy compact according to claim 13,wherein the suffixes x, y, z, w, and v by atomic percent satisfy80≦x≦92, 1≦y≦5, 3≦z≦10, 3≦w≦7, and 0≦v≦5, respectively.
 15. A hardmagnetic alloy compact according to claim 1, wherein said hard magneticalloy compact has the following formula: T_(x)M_(y)R_(z)B_(w)E_(v)Si_(u)wherein T represents at least one element selected from the groupconsisting of Fe, Co and Ni, M represents at least one element selectedfrom the group consisting of Zr, Nb, Ta and Hf, R represents at leastone rare earth element, E represents at least one element selected fromthe group consisting of Cr, Al, Pt and platinum elements, and thesuffixes x, y, z, w, v, and u by atomic percent satisfy 50≦x, 0≦y≦15,3≦z≦20, 2≦w≦20, 0≦v≦10, and 0≦u≦5, respectively.
 16. A hard magneticalloy compact according to claim 15, wherein the suffixes x, y, z, w, v,and u by atomic percent satisfy 80≦x≦92, 1≦y≦5, 3≦z≦10, 3≦w≦7, 0≦v≦5,and 0.5≦u≦5, respectively.
 17. A method for producing a hard magneticalloy compact comprising: quenching one of an Fe- and FeCo-based alloycontaining 3 to 20 atomic percent of at least one rare earth element Rand 2 to 20 atomic percent of B to form a texture, in which at least apart of the texture comprises one of an amorphous phase and finecrystalline phase having an average crystal grain size of not greaterthan 100 nm; performing one of crystallization and grain growth understress, such that a mixed phase composed of a hard magnetic phase andone of a soft magnetic and semi-hard magnetic phase is formed in saidtexture; and imparting anisotropy to a crystal axis of said hardmagnetic phase.
 18. A method for producing a hard magnetic alloy compactaccording to claim 17, wherein after performing one of crystallizationand grain growth of said alloy under stress, said alloy is annealed at400° C. to 1000° C. to precipitate a fine crystalline phase having anaverage crystal grain size of not greater than 100 nm as a main phase inthe texture.
 19. A method for producing a hard magnetic alloy compactaccording to claim 17, wherein said alloy after quenching is compactedwhile performing one of crystallization and grain growth of said alloyunder stress.
 20. A method for producing a hard magnetic alloy compactaccording to claim 17, wherein said hard magnetic alloy contains anamorphous phase, and is formed by solidifying an alloy having hardmagnetic characteristics when being crystallized, by means of asoftening phenomenon of said alloy during the crystallization reaction.21. A method for producing a hard magnetic alloy compact according toclaim 17, wherein said alloy is heated under stress.